JPWO2016190370A1 - Steel sheet and manufacturing method thereof - Google Patents

Abstract

Low-carbon steel sheet with excellent cold forgeability and impact resistance after carburizing, quenching, and tempering, having a predetermined component composition, an average particle size of carbide of 0.4 μm to 2.0 μm, and a pearlite area ratio 6% or less, the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains is more than 1, and the Vickers hardness is 100HV or more and 180HV or less.

Description

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

  For steel sheets containing 0.1 to 0.4% carbon by mass%, from blank materials, press forming, hole expanding forming, bending forming, drawing forming, thickening and thinning forming, and combinations thereof Cold forging is applied, and it is used as a material for driving system parts such as automobile gears and clutches. Compared to conventional hot forging and the like, cold forging increases the amount of strain accumulated in the material, causing cracks in the material and buckling at the time of molding, causing a problem of deterioration of component characteristics.

  In particular, after carburizing and tempering the molding material in order to obtain wear resistance, residual stress is generated by the heat treatment, so that cracking and buckling are caused and the cracks are generated and propagated. In order to use it as a drive system component, it is required to acquire impact resistance characteristics so as not to break brittlely against a heavy load applied instantaneously due to gear engagement at the start. Steel sheets are required to ensure excellent cold forgeability and impact resistance after carburizing, quenching and tempering.

  So far, many proposals have been made on techniques for improving the cold forgeability of steel sheets and the impact resistance after carburizing (see, for example, Patent Documents 1 to 5).

  For example, in Patent Document 1, as steel for mechanical structure whose toughness is improved by suppressing coarsening of crystal grains in carburizing heat treatment, C: 0.10 to 0.30%, Si: 0.05 -2.0%, Mn: 0.10 to 0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 to 3.00%, Al: 0.005 0.050%, Nb: 0.02 to 0.10%, N: 0.0300% or less, consisting of the balance Fe and inevitable impurities, the structure before cold working is a ferrite pearlite structure, A steel for machine structural use having an average ferrite grain size of 15 μm or more is disclosed.

  In Patent Document 2, as steel having excellent cold workability and carburizing hardenability, C: 0.15 to 0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, B: 0.005 to 0.050%, the balance is composed of Fe and unavoidable impurities, and mainly comprises a ferrite phase and a graphite phase. A steel having a structure that achieves this is disclosed.

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

  In Patent Document 4, cold forging is performed after spheroidizing annealing, and the parts produced in the carburizing and quenching and tempering process have excellent workability while suppressing the coarsening of crystal grains even in subsequent carburizing. However, steel for carburized parts having excellent impact resistance characteristics and impact fatigue resistance characteristics is disclosed.

  In Patent Document 5, as cold tool steel for plasma carburizing, C: 0.40-0.80%, Si: 0.05-1.50%, Mn: 0.05-1.50%, and V: 1.8 to 6.0%, Ni: 0.10 to 2.50%, Cr: 0.1 to 2.0%, and Mo: 3.0% or less Or the steel which contains 2 or more types and remainder consists of Fe and an unavoidable impurity is disclosed.

JP2013-040376A Japanese Patent Laid-Open No. 06-116679 JP 09-201644 A JP 2006-213951 A JP-A-10-158780

  The structure of the steel for machine structural use in Patent Document 1 is a structure of ferrite + pearlite, which has a larger hardness than the structure of ferrite + cementite, and therefore suppresses wear of the mold in cold forging. It cannot be said that it is necessarily a steel for machine structural use with excellent cold forgeability.

  In the steel of Patent Document 2, annealing at a high temperature is essential for the cementitization treatment of cementite, and it is impossible to suppress a decrease in yield and an increase in manufacturing cost.

  The manufacturing method of Patent Document 3 requires hot forging after cold forging and carburizing, and since hot forging is essential, it is not a manufacturing method that leads to drastic cost reduction.

  It is unclear whether the steel for carburized parts of Patent Document 4 can achieve the same effect in cold forging where a large strain is applied, and there is no specific structure form or structure control method. Since it is clear, it cannot be said that the steel shows excellent workability even in the forming forging by applying a large strain in the cold such as plate forging which has been widely applied in recent years.

  Patent Document 5 does not disclose any knowledge and technique relating to optimum components and microstructures for improving the formability of steel, particularly cold forgeability.

  The present invention is a steel plate suitable for obtaining parts such as a high cycle gear by sheet forming, and a method for producing the steel plate, which are excellent in cold forgeability and impact resistance after carburizing and quenching and tempering, in view of the above-described prior art. The issue is to provide.

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

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

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

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

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

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

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

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

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

  (1) Component composition is mass%, C: 0.10-0.40%, Si: 0.01-0.30%, Mn: 0.30-1.00%, Al: 0.001- 0.10%, Cr: 0.50-2.00%, Mo: 0.001-1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020% or less O: 0.020% or less, Ti: 0.010% or less, B: 0.0005% or less, Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or less, Nb : 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0 0.050% or less, Ca: 0.050% or less, Y: 0.050% or less, Zr: 0.050% or less, La: 0.050% or less, and Ce: 0.05% or less. A low carbon steel plate that contains 50%, the balance being Fe and impurities, wherein the metal structure of the low carbon steel plate has a carbide grain size of 0.4 to 2.0 μm, a pearlite area ratio of 6% or less, and ferrite grains A steel sheet characterized in that the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the steel satisfies more than 1, and the Vickers hardness of the low carbon steel sheet is 100HV or more and 180HV or less.

  (2) A manufacturing method for manufacturing the steel sheet of (1), wherein the steel slab having the component composition of (1) is hot-rolled to complete finish hot rolling in a temperature range of 650 ° C. to 950 ° C. The hot-rolled steel sheet is wound up at 400 ° C. or higher and 600 ° C. or lower, and the wound hot-rolled steel sheet is pickled, and the pickled hot-rolled steel sheet is 30 ° C./hour or higher and 150 ° C./hour or lower. At a heating rate of 650 ° C. or more and 720 ° C. or less, the first stage annealing is performed for 3 hours or more and 60 hours or less, and then the hot-rolled steel sheet is heated at 1 ° C./hour or more and 80 ° C./hour. Heating to an annealing temperature of 725 ° C. or higher and 790 ° C. or lower at the following heating rate and performing second-stage annealing for holding for 3 hours or longer and 50 hours or shorter, Cooling to 650 ° C. at a cooling rate of ℃ / hour or less Manufacturing method of forging property and steel plate.

  ADVANTAGE OF THE INVENTION According to this invention, the steel plate suitable for obtaining components, such as a high cycle gear, can be provided, which is excellent in cold forgeability and impact resistance characteristics after carburizing and quenching and tempering, and particularly by plate forming.

It is a figure which shows typically the aspect of the crack introduced by the outline | summary of a cold forging test, and cold forging. (A) shows a disk-shaped test material cut out from a hot-rolled steel sheet, (b) shows the shape of the test material after cold forging, and (c) shows a cross-sectional aspect of the test material after cold forging. Indicates. It is a figure which shows typically the outline | summary of the drop weight test which evaluates the impact resistance characteristic of the sample which performed carburizing quenching tempering. It is a figure which shows the relationship between the ratio of the number of intergranular carbide | carbonized_materials with respect to the number of intragranular carbide | carbonized_material, the crack length of a cold forging test piece, and the impact resistance property after carburizing quenching and tempering. It is a figure which shows another relationship between the ratio of the number of intergranular carbide | carbonized_materials with respect to the number of intragranular carbide | carbonized_materials, the crack length of a cold forging test piece, and the impact resistance property after carburizing quenching and tempering.

  Hereinafter, the present invention will be described in detail. First, the reasons for limiting the component composition of the steel sheet of the present invention will be described. Here, “%” related to the component composition means “mass%”.

[C: 0.10 to 0.40%]
C is an element that forms carbides in steel and is effective in strengthening steel and refining ferrite grains. In order to suppress the occurrence of satin in cold working and ensure the surface aesthetics of the cold forged parts, it is essential to suppress the coarsening of the ferrite grain size, but if it is less than 0.10%, the volume fraction of carbides Is insufficient, and it becomes impossible to suppress the coarsening of the carbide during annealing, 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 instantaneously applied, a large amount of cracks that become the starting point of fracture are generated, leading to a decrease in impact resistance. 0.40% or less. Preferably it is 0.38% or less.

[Si: 0.01-0.30%]
Si is an element that acts as a deoxidizer and affects the morphology of carbides. In order to reduce the number of carbides in the ferrite grains to obtain the deoxidation effect and increase the number of carbides on the ferrite grain boundaries, the austenite phase is generated during annealing by the two-step type annealing, It is necessary to dissolve the carbide and then slowly cool it to promote the formation of carbide at the ferrite grain boundaries.

  If Si exceeds 0.30%, the ductility of ferrite is reduced, cracking is likely to occur during cold forging, and cold forgeability and impact resistance after carburizing and quenching are reduced. % Or less. Preferably it is 0.28% or less.

  Si is preferably as small as possible, but a reduction to less than 0.01% causes a significant increase in refining costs, so Si is set to 0.01% or more. Preferably it is 0.02% or more.

[Mn: 0.30 to 1.00%]
Mn is an element that controls the form of carbide in the two-step annealing. If it is less than 0.30%, it becomes difficult to form carbides on the ferrite grain boundaries in the slow cooling after the second stage annealing, so Mn is 0.30% or more. Preferably it is 0.33% or more.

  On the other hand, if it exceeds 1.00%, the toughness after carburizing, quenching, and tempering decreases, so Mn is made 1.00% or less. Preferably it is 0.96% or less.

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

  On the other hand, if it exceeds 0.10%, the number ratio of carbides on the grain boundary is lowered and the crack length at the time of cold forging is increased, so Al is made 0.10% or less. Preferably it is 0.09% or less.

[Cr: 0.50 to 2.00%]
Cr and Mo are elements that improve toughness. Cr is an element effective for stabilizing carbide during heat treatment. If it is less than 0.50%, it becomes difficult to leave carbides during carburization, leading to a coarsening of the austenite grain size in the surface layer and a reduction in impact resistance, so Cr is made 0.50% or more. Preferably it is 0.52% or more.

  On the other hand, if it exceeds 2.00%, the concentration of Cr in the carbide increases, and a lot of fine carbides remain in the austenite phase produced by the two-step annealing, so that the grains after slow cooling Since carbides are also present therein, the increase in hardness and the number ratio of grain boundary carbides are reduced, and the cold forgeability is reduced, so Cr is made 2.00% or less. Preferably it is 1.94% or less.

[Mo: 0.001 to 1.00%]
Mo is an element effective for controlling the form of carbide. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Mo is made 0.001% or more. Preferably it is 0.017% or more.

  On the other hand, if it exceeds 1.00%, Mo concentrates in the carbide, and stable carbide increases in the austenite phase. Therefore, carbide is also present in the grains after slow cooling, and the increase in hardness and grain boundary carbides. Since the number ratio is lowered and the cold forgeability is lowered, Mo is made 1.00% or less. Preferably it is 0.94% or less.

  The following elements are impurities and need to be controlled below a certain amount.

[P: 0.020% or less]
P is an element that segregates at the ferrite grain boundaries and suppresses the formation of grain boundary carbides. The smaller the number, the better. The P content may be 0, but in order to achieve a purity of less than 0.0001% in the refining process, refining takes a long time and causes a significant increase in manufacturing costs. 0.0001 to 0.0013%.

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

[S: 0.010% or less]
S is an impurity element that forms non-metallic inclusions such as MnS. Non-metallic inclusions are the starting point of cracking during cold forging, so the smaller the S, the better. The S content may be 0, but if S is reduced to less than 0.0001%, the refining cost is greatly increased, so the practical lower limit is 0.0001 to 0.0012%.

  On the other hand, if it exceeds 0.010%, the crack length during cold forging is increased, so S is made 0.010% or less. Preferably it is 0.009% or less.

[N: 0.020% or less]
N is an element that segregates to the ferrite grain boundary and suppresses the formation of carbides on the grain boundary. The smaller the number, the better. The N content may be 0, but if the content is reduced to less than 0.0001%, the refining cost is greatly increased, so the substantial lower limit is 0.0001 to 0.0006%.

  On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains is less than 1 even when two-phase annealing and annealing are performed, and the cold forgeability is low. Since N falls, 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 forms an oxide in steel. Since the oxide which exists in a ferrite grain turns into a production | generation site | part of a carbide | carbonized_material, the one where it is smaller is preferable. The content of O may be 0, but if O is reduced to less than 0.0001%, the refining cost is greatly increased, so the practical lower limit is 0.0001 to 0.0006%.

  On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains becomes less than 1, and cold forgeability decreases, so O is 0.020% or less. To do. Preferably it is 0.017% or less.

[Ti: 0.010% or less]
Ti is an element that is important for controlling the form of carbides, and is an element that promotes the formation of carbides in ferrite grains when contained in a large amount. The Ti content may be 0, but if the content is reduced to less than 0.0001%, the refining cost increases significantly, so the substantial lower limit is 0.0001 to 0.0006%.

  On the other hand, if it exceeds 0.010%, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains becomes less than 1, and the cold forgeability decreases, so Ti is 0.010% or less. To do. Preferably it is 0.007% or less.

[B: 0.0005% or less]
B is an element effective for controlling dislocation slip during cold forging. Since the activity of the slip system is limited by the inclusion of a large amount, B is preferably as small as possible. The B content may be zero. Careful attention is required for the detection of B less than 0.0001%, and depending on the analyzer, the detection is below the lower limit of detection.

  On the other hand, if it exceeds 0.0005%, dislocation cross-slip is suppressed in the shear band formed by cold forging, and strain is concentrated locally to cause cracking. Therefore, B is set to 0.0005% or less. . Preferably it is 0.0005% or less.

[Sn: 0.050% or less]
Sn is an element mixed in from the steel raw material (scrap), and the smaller the better. The Sn content may be 0, but if the content is reduced to less than 0.001%, the refining cost increases significantly, so the substantial lower limit is 0.001 to 0.002%.

  On the other hand, if it exceeds 0.050%, ferrite becomes brittle and cold forgeability deteriorates, so Sn is made 0.050% or less. Preferably it is 0.048% or less.

[Sb: 0.050% or less]
Sb is an element mixed from steel raw material (scrap) like Sn. Since Sb segregates at the grain boundaries and reduces the number ratio of grain boundary carbides, the smaller the Sb, the better. The Sb content may be 0, but if the content is reduced to less than 0.001%, the refining cost is greatly increased, so the substantial lower limit is 0.001 to 0.002%.

  On the other hand, if it exceeds 0.050%, the cold forgeability deteriorates, so Sb is made 0.050% or less. Preferably it is 0.048% or less.

[As: 0.050% or less]
As is an element mixed from steel raw material (scrap), as is Sn and Sb, As is segregated at the grain boundary and lowers the number ratio of grain boundary carbides. The content of As may be 0, but if the content is reduced to less than 0.001%, the refining cost is greatly increased, so the substantial lower limit is 0.001 to 0.002%.

  On the other hand, if it exceeds 0.050%, the number ratio of grain boundary carbides decreases and cold forgeability decreases, so As is made 0.050% or less. Preferably it is 0.045% or less.

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

[Nb: 0.10% or less]
Nb is an element effective for controlling the form 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 sufficiently obtained, so Nb is preferably 0.001% or more. More preferably, it is 0.002% or more.

  On the other hand, if it exceeds 0.10%, a large number of fine Nb carbides precipitate, the strength excessively increases, the number ratio of grain boundary carbides decreases, and cold forgeability decreases, so Nb is 0 10% or less. Preferably it is 0.09% or less.

[V: 0.10% or less]
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 sufficiently obtained, so V is preferably 0.001% or more. More preferably, it is 0.004% or more.

  On the other hand, if it exceeds 0.10%, a large number of fine V carbides precipitate, the strength excessively increases, the number ratio of grain boundary carbides decreases, and the cold forgeability decreases, so V is 0. 10% or less. Preferably it is 0.09% or less.

[Cu: 0.10% or less]
Cu is an element that forms fine precipitates and contributes to improvement in strength. If it is less than 0.001%, the effect of improving the strength cannot be obtained sufficiently, so Cu is preferably made 0.001% or more. More preferably, it is 0.008% or more.

  On the other hand, if it exceeds 0.10%, red hot brittleness will develop during hot rolling and the productivity will decrease, so Cu is made 0.10% or less. Preferably it is 0.09% or less.

[W: 0.10% or less]
W, like Nb and V, is an element effective for controlling the form of carbide. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so W is preferably 0.001% or more. More preferably, it is 0.003% or more.

  On the other hand, if it exceeds 0.10%, many fine W carbides precipitate, the strength excessively increases, the number ratio of grain boundary carbides decreases, and the cold forgeability decreases, so W is 0. 10% or less. Preferably it is 0.08% or less.

[Ta: 0.10% or less]
Ta, as well as Nb, V, and W, is an element effective for controlling the morphology of carbides. If less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ta is preferably 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 precipitate, the strength increases excessively, the number ratio of grain boundary carbides decreases, and the cold forgeability decreases, so Ta is 0. 10% or less. Preferably it is 0.09% or less.

[Ni: 0.10% or less]
Ni is an element effective for improving the impact resistance characteristics of parts. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ni is preferably made 0.001% or more. More preferably, it is 0.002% or more.

  On the other hand, if it exceeds 0.10%, the number ratio of grain boundary carbides decreases and cold forgeability decreases, so Ni is made 0.10% or less. Preferably it is 0.09% or less.

[Mg: 0.050% or less]
Mg is an element that can control the form of sulfide by addition of a small amount. If it is less than 0.0001%, the effect of addition cannot be sufficiently obtained, so Mg is preferably made 0.0001% or more. More preferably, it is 0.0008% or more.

  On the other hand, if it exceeds 0.050%, ferrite becomes brittle and cold forgeability deteriorates, so Mg is made 0.050% or less. Preferably it is 0.049% or less.

[Ca: 0.050% or less]
Ca, like Mg, is an element that can control the form of sulfide with a small amount of addition. If less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ca is preferably made 0.001% or more. More preferably, it is 0.003% or more.

  On the other hand, if it exceeds 0.050%, coarse Ca oxide is generated and becomes a starting point of crack generation during cold forging, so Ca is made 0.050% or less. Preferably it is 0.04% or less.

[Y: 0.050% or less]
Y, like Mg and Ca, is an element that can control the form of sulfide by addition of a trace amount. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Y is preferably 0.001% or more. More preferably, it is 0.003% or more.

  On the other hand, if it exceeds 0.050%, coarse Y oxide is generated and becomes a starting point of cracking during cold forging, so Y is made 0.050% or less. Preferably it is 0.031% or less.

[Zr: 0.050% or less]
Zr, like Mg, Ca, and Y, is an element that can control the form of sulfide by adding a small amount. If it is less than 0.001%, the effect of addition cannot be obtained sufficiently, so Zr is preferably 0.001% or more. More preferably, it is 0.004% or more.

  On the other hand, if it exceeds 0.050%, a coarse Zr oxide is generated and becomes a starting point of cracking during cold forging, so Zr is made 0.050% or less. Preferably it is 0.045% or less.

[La: 0.050% or less]
La is an element that is effective in controlling the form of sulfides when added in a small amount, and is an element that segregates at the grain boundaries and lowers the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect cannot be obtained sufficiently, so La is preferably 0.001% or more. More preferably, it is 0.003% or more.

  On the other hand, if it exceeds 0.050%, the number ratio of grain boundary carbides decreases and cold forgeability decreases, so La is made 0.050% or less. Preferably it is 0.047% or less.

[Ce: 0.050% or less]
Ce, like La, is an element that can control the form of sulfide by addition of a small amount, and is an element that segregates at the grain boundary and lowers the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect cannot be obtained sufficiently, so Ce is preferably 0.001% or more. More preferably, it is 0.003% or more.

  On the other hand, if it exceeds 0.050%, the number ratio of grain boundary carbides decreases and cold forgeability decreases, so Ce is made 0.050% or less. Preferably it is 0.046% or less.

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

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

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

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

  In cold forging, strong processing exceeding 1 is performed with considerable strain. For this reason, in the conventional steel plate, generation of voids and / or cracks accompanying an increase in dislocation density cannot be prevented, and it has been difficult to improve cold forgeability.

  In order to solve this difficult problem, it is effective to suppress the formation of shear bands during molding. From the viewpoint of the microstructure, the formation of a shear band can be understood as a phenomenon in which a slip generated in one grain overcomes the grain boundary and continuously propagates to adjacent grains. Therefore, in order to suppress the formation of shear bands, it is necessary to prevent the propagation of slip across the crystal grain boundary.

  The carbide in the steel plate is a strong particle that prevents slipping, and by allowing the carbide to exist at the ferrite grain boundary, formation of a shear band can be suppressed and cold forgeability can be improved.

  In order to obtain such an effect, the carbide must be dispersed in an appropriate size in the metal structure. Therefore, the average particle diameter of the carbide is set to 0.4 μm or more and 2.0 μm or less. When the particle diameter of the carbide is less than 0.4 μm, the hardness of the steel sheet is remarkably increased and the cold forgeability is lowered. More preferably, it is 0.6 μm or more.

  On the other hand, if the average particle diameter of the carbide exceeds 2.0 μm, the carbide becomes a starting point of cracking during cold forming. More preferably, it is 1.95 μm or less.

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

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

  Based on the theory and principle, cold forgeability is considered to be strongly influenced by the carbide coverage of ferrite grain boundaries, and its high-precision measurement is required, but carbide to ferrite grain boundaries in a three-dimensional space. In order to measure the coverage ratio, serial sectioning SEM observation or three-dimensional EBSP observation in which the cutting and observation of the sample by FIB is repeatedly performed in a scanning electron microscope is indispensable. Accumulation is essential.

  The present inventors have clarified this, and as a result of searching for a simpler and more accurate evaluation index, if the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains is used as an index, the cold The present inventors have found that forgeability can be evaluated, and if the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains exceeds 1, the cold forgeability is significantly improved. .

  In addition, since any buckling, folding, and folding of the steel sheet that occurs during cold forging are caused by strain localization accompanying the formation of a shear band, similarly, carbide should be present at the ferrite grain boundaries. Thus, if the formation of shear bands and the localization of strain are alleviated, the occurrence of buckling, folding, and folding can be suppressed.

  Carbide is observed with a scanning electron microscope. Prior to observation, a sample for tissue observation was polished with emery paper by wet polishing and diamond abrasive grains having an average particle size of 1 μm, and the observation surface was finished to a mirror surface, and the tissue was then saturated with a saturated picric acid-alcohol solution. Etch.

  The observation magnification is set to 3000 times, and eight images of a 30 μm × 40 μm visual field in a 1/4 layer thickness are taken at random. The area of each carbide contained in the region is measured in detail with the image analysis software represented by Mitani Corporation (Win ROOF) on the obtained tissue image. The equivalent circle diameter (= 2 × √ (area / 3.14)) is determined from the area of each carbide, and the average value is defined as the carbide particle diameter.

In order to suppress the influence of measurement error due to noise, carbides having an area of 0.01 μm ² or less are excluded from the evaluation target.

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

  Cold forgeability can be improved by setting the ferrite grain size to 3.0 μm or more and 50.0 μm or less as the structure after annealing. If the ferrite particle size is less than 3 μm, the hardness increases and cracks and cracks are likely to occur during cold forging, so the ferrite particle size is preferably 3.0 μm or more. More preferably, it is 7.5 μm or more.

  On the other hand, if the ferrite grain size exceeds 50.0 μm, the number of carbides on the grain boundaries that suppress the propagation of slip is reduced and the cold forgeability is lowered. Therefore, the ferrite grain size is preferably 50.0 μm or less. . More preferably, it is 37.9 μm or less.

  The ferrite grain size was determined by observing the structure of the observation surface etched with a 3% nitric acid-alcohol solution with an optical microscope or a scanning electron microscope after polishing the observation surface of the sample for tissue observation to a mirror surface in the above procedure. The line segment method is applied to the captured image.

  By setting the Vickers hardness of the steel sheet to 100 HV or more and 180 HV or less, it is possible to improve cold forgeability and impact resistance characteristics after carburizing and tempering. If the Vickers hardness is less than 100 HV, buckling is likely to occur during cold forging, and folding and folding of the buckled portion occur and impact resistance is reduced. Therefore, the Vickers hardness is 100 HV or more. To do. Preferably it is 110HV or more.

  On the other hand, if the Vickers hardness exceeds 180 HV, the ductility is lowered, internal cracks are liable to occur during cold forging, and the impact resistance is deteriorated, so the Vickers hardness is set to 180 HV or less. Preferably it is 170 HV or less.

  Next, a method for evaluating cold forgeability will be described.

  FIG. 1 schematically shows the outline of the cold forging test and the mode of cracks introduced by cold forging. FIG. 1 (a) shows a disk-shaped test material cut out from a hot-rolled steel sheet, FIG. 1 (b) shows the shape of the test material after cold forging, and FIG. 1 (c) shows the result after cold forging. The cross-sectional aspect of this test material is shown.

  As shown in FIG. 1, a disk-shaped test material 1 having a diameter of 70 mm is cut out from a hot-rolled steel sheet having a thickness of 5.2 mm (see FIG. 1 (a)). A test material is prepared (not shown). Next, using a Mori Tekko one-shot forming press machine, the vertical wall of the cup-shaped test material is thickened (cold forging) at a thickness increase ratio of 1.54 (= 8 mm / 5.2 mm). A cup-shaped test material 2 having a diameter of 30 mm, a height of 30 mm, and a vertical wall thickness of 8 mm is produced (see FIG. 1B).

  The cup-shaped test material 2 subjected to the thickening molding is cut with a wire cut electric discharge machine manufactured by FANUC so that the cross section of the diameter portion appears (see FIG. 1C). The cut surface is mirror-polished to confirm that the crack 3 is present on the cut surface, and the ratio of the maximum length L of the crack existing in the vertical wall portion to the thickness of the vertical wall portion after thickening (= maximum of the crack) Length L / thickness of the vertical wall after thickening is 8 mm). The cold forgeability is evaluated by this measured value.

  In addition, even when the initial plate thickness is other than 5.2 mm, the diameter of the disc-shaped test material to be cut out is adjusted so that the height of the vertical wall after the thickness increase becomes 30 mm, and the same thickness increase ratio of 1.54 If it shape | molds, since an evaluation result can be reproduced irrespective of initial stage plate | board thickness, the hot-rolled steel plate which this invention makes object is not limited to the hot-rolled steel plate of plate | board thickness 5.2mm. The present invention can improve cold forgeability and impact resistance after carburizing and tempering even in a hot rolled steel sheet having a general thickness (2 to 15 mm).

  Next, the manufacturing method of the present invention will be described. The technical idea of the manufacturing method of the present invention is to consistently manage hot rolling conditions and annealing conditions when manufacturing steel sheets from steel slabs having the above-mentioned composition, and to provide cold forgeability and impact resistance characteristics after carburizing and quenching and tempering. It is to improve.

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

[Characteristics of hot rolling]
A molten steel having a required composition is continuously cast to form a slab, and the slab is subjected to hot rolling as usual, or after being cooled and heated, and then subjected to hot rolling, 650 ° C. to 950 ° C. Finish hot rolling in the following temperature range. The hot-rolled steel sheet after finish rolling is cooled on the ROT and wound at a winding temperature of 400 ° C. or higher and 600 ° C. or lower.

[Characteristics of annealing]
The hot-rolled steel sheet is subjected to two-step annealing that is held in two temperature ranges after pickling. At that time, in the first-stage annealing, the hot-rolled steel sheet is heated to an annealing temperature of 30 ° C./hour or more to 150 ° C./hour or more. Heating is performed at a heating rate of ℃ / hour or less, and annealing is performed in a temperature range of 650 ° C. to 720 ° C. for 3 hours to 60 hours.

  In the next second stage annealing, the hot-rolled steel sheet is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour, and in a temperature range of 725 ° C. to 790 ° C. for 3 hours to 50 hours. The following annealing is performed.

  Next, 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.

  A low carbon steel sheet excellent in cold forgeability and impact resistance after carburizing and quenching and tempering can be obtained by cooperation between the hot rolling conditions and the annealing conditions.

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

[Hot rolling]
Finishing hot rolling temperature: 650 ° C or higher and 950 ° C or lower Winding temperature: 400 ° C or higher and 600 ° C or lower

  The molten steel having the required composition is continuously cast into a slab, as it is, or once cooled and heated, and then subjected to hot rolling, and finish hot rolling is finished at a temperature range of 650 ° C. to 950 ° C. The rolled steel sheet is wound up at 400 ° C. or higher and 600 ° C. or lower.

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

  If the heating temperature exceeds 1300 ° C or the heating time exceeds 7 hours, decarburization of the slab surface layer becomes prominent, and the austenite grains on the surface layer grow abnormally during heating before quenching, resulting in reduced impact resistance. Therefore, 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.

  The finish hot rolling ends at a temperature of 650 ° C. or higher and 950 ° C. or lower. If the finishing hot rolling temperature is less than 650 ° C., the rolling load is remarkably increased due to an increase in deformation resistance of the steel material, and further, the roll wear amount is increased and the productivity is lowered. Therefore, the finishing hot rolling temperature is 650 ° C. That's it. Preferably it is 680 degreeC or more.

  On the other hand, when the finish hot rolling temperature exceeds 950 ° C., a thick scale is generated while passing through the ROT (Run Out Table), and the surface of the steel sheet is wrinkled due to the scale, during cold forging, and / or Alternatively, when an impact load is applied after carburizing, quenching, and tempering, cracks are generated starting from the flaws and impact resistance is reduced, so the finish hot rolling temperature is 950 ° C. or lower. Preferably it is 920 degrees C or less.

  The cooling rate when the hot-rolled steel sheet is cooled on the ROT is preferably 10 ° C./second or more and 100 ° C./second or less. When the cooling rate is less than 10 ° C./second, the generation of thick scale and the generation of wrinkles due to it cannot be suppressed during the cooling, and the impact resistance is reduced. Seconds or more are preferred. More preferably, it is 20 ° C./second or more.

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

  When the hot-rolled steel sheet at 100 ° C. to room temperature is taken out after winding, microcracks are generated in the low-temperature transformation structure, and it is difficult to remove the cracks in the subsequent pickling process and cold rolling process. Alternatively, when an impact load is applied after carburizing, quenching, and tempering, the crack progresses starting from the crack and causes a reduction in impact resistance. Therefore, the cooling rate is preferably 100 ° C./second or less. More preferably, it is 80 ° C./second or less.

  Note that the cooling rate is as follows: from the time when the hot-rolled steel sheet after finish hot rolling passes through the non-water-injection section to the time when the water-cooling in the water-injection section is cooled to the winding target temperature on the ROT It refers to the cooling capacity received from the cooling equipment in the water injection section, and does not indicate the average cooling rate from the water injection start point to the temperature taken up by the winder.

  The coiling temperature is 400 ° C. or higher and 600 ° C. or lower. This is a temperature lower than the general winding temperature. By winding the hot-rolled steel sheet manufactured under the above-described conditions in this temperature range, the structure of the steel sheet can be a bainite structure in which carbides are dispersed in fine ferrite.

  When the coiling temperature is less than 400 ° C., austenite that has not been transformed before winding is transformed into hard martensite, and when the wound hot-rolled steel sheet is discharged, cracks are generated in the surface layer, resulting in impact resistance. Since the temperature falls, the coiling temperature is set to 400 ° C. or higher. Preferably it is 430 degreeC or more.

  On the other hand, when the coiling temperature exceeds 600 ° C., pearlite with a large lamellar spacing is generated, and thick needle-like carbides having high thermal stability are formed. Even after the two-step annealing, the needle-like carbides are formed. Remains. At the time of cold forging, cracks are generated and propagated starting from these needle-like carbides, so the coiling temperature is set to 600 ° C. or lower. Preferably it is 570 degrees C or less.

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

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

  By performing the first stage annealing in a temperature range of Ac1 point or less, the carbide is coarsened and the added gold element is concentrated to enhance the thermal stability of the carbide. Thereafter, the temperature is raised to a temperature range of Ac1 point or higher, austenite is generated in the structure, carbides in fine ferrite grains are dissolved in austenite, and coarse carbides remain in the austenite.

  By subsequent slow cooling, austenite is transformed into ferrite, and the carbon concentration in the austenite is increased. By gradually cooling, carbon atoms are adsorbed on the carbide remaining in the austenite, and the carbide and austenite cover the ferrite grain boundaries, and finally a structure in which a large amount of carbide exists in the ferrite grain boundaries. Can be formed. Therefore, it is clear that the structure defined in the present invention cannot be formed only by simple annealing.

  Specific annealing conditions will be described below.

[First stage annealing]
Heating rate to annealing temperature: 30 ° C / hour or more and 150 ° C / hour Annealing temperature: 650 ° C or more and 720 ° C or less Holding time at annealing temperature: 3 hours or more and 60 hours or less

  The heating rate up to the first stage annealing temperature is 30 ° C./hour or more and 150 ° C./hour or less. If the heating rate is less than 30 ° C./hour, it takes time to raise the temperature and the productivity decreases, so the heating rate is set to 30 ° C./hour or more. Preferably, it is 40 ° C./hour or more.

  On the other hand, when the heating rate exceeds 150 ° C./hour, the temperature difference between the outer peripheral portion and the inside of the coil increases, so that frost and seizure occur due to the difference in thermal expansion, and irregularities are generated on the surface of the steel sheet. At the time of cold forging, cracks are generated starting from the unevenness, which causes a decrease in cold forgeability and a decrease in impact resistance after carburizing and quenching and tempering. Therefore, the heating rate is set to 150 ° C./hour or less. Preferably it is 120 degrees C / hour or less.

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

  On the other hand, if the annealing temperature exceeds 720 ° C., austenite is generated before the stability of the carbide is increased, and the above-described change in structure cannot be controlled. Therefore, the annealing temperature is set to 720 ° C. or less. Preferably it is 700 degrees C or less.

  The holding time in the first stage annealing (first stage holding time) is 3 hours or more and 60 hours or less. If the first stage holding time is less than 3 hours, the carbide is not sufficiently stabilized, and it is difficult to leave the carbides in the second stage annealing, so the first stage holding time is 3 hours. That's it. Preferably it is 10 hours or more.

  On the other hand, if the first stage holding time exceeds 60 hours, no further improvement in the stability of the carbide can be expected, and the productivity is further lowered. Therefore, the first stage holding time is set to 60 hours or less. Preferably it is 50 hours or less.

[Second stage annealing]
Heating rate to annealing temperature: 1 ° C / hour or more and 80 ° C / hour Annealing temperature: 725 ° C or more and 790 ° C or less Holding time at annealing temperature: 3 hours or more and 50 hours or less

  After completion of the holding in the first stage annealing, the hot-rolled steel sheet is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour. When cooled without performing the second-stage annealing, the ferrite grain size does not increase and an ideal structure cannot be obtained.

  In the second stage annealing, austenite is generated and grows from the ferrite grain boundary. By slowing the heating rate, nucleation of austenite can be suppressed, and the grain boundary coverage of carbides can be increased in the structure obtained after slow cooling. Therefore, it is preferable that the heating rate in the second stage annealing is small.

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

  On the other hand, if the heating rate exceeds 80 ° C./hour, the temperature difference between the outer peripheral portion and the inside of the coil increases, resulting in a large difference in thermal expansion due to transformation, causing scouring and seizure, and unevenness on the steel sheet surface. Produces. At the time of cold forging, cracks are generated starting from this unevenness, which causes a decrease in cold forgeability and a decrease in impact resistance after carburizing and quenching and tempering, so the heating rate is 80 ° C./hour or less.

  The annealing temperature in the second stage annealing (second stage annealing temperature) is 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 reduced, and after cooling after the second stage annealing, the number ratio of carbides on the ferrite grain boundaries is reduced. Becomes smaller. Therefore, 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 annealing temperature of the second stage exceeds 790 ° C, it becomes difficult to leave the carbide in the austenite, and it becomes difficult to control the above-described structural change, so the annealing temperature of the second stage is 790 ° C or less. And Preferably it is 780 degrees C or less.

  The holding time in the second stage annealing (second stage holding time) is 1 hour to 50 hours. If the second stage holding time is less than 1 hour, the amount of austenite produced is small, and the carbides in the ferrite grains are not sufficiently dissolved, and the number ratio of carbides on the ferrite grain boundaries can be increased. Since it becomes difficult and the ferrite grain size becomes small, the second stage holding time is set to 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 the carbide in the austenite. Therefore, the second stage holding time is set to 50 hours or less. Preferably it is 45 hours or less.

[Cooling after annealing]
Cooling stop temperature: 650 ° C
Cooling rate: 1 ° C / hour or more and 100 ° C / hour or less

  After the holding in the second stage annealing is completed, the annealed hot-rolled steel sheet is gradually cooled to 650 ° C. at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less. When the slow cooling stop temperature exceeds 650 ° C., the untransformed austenite is transformed into pearlite or bainite by the cooling rate exceeding 100 ° C./hour until the room temperature, the hardness increases, and the cold forgeability increases. Since it decreases, the cooling stop temperature is set to 650 ° C.

  In order to cool the austenite produced in the second stage annealing to transform it into ferrite and to adsorb carbon to the carbide remaining in the austenite, it is preferable that the cooling rate is slow. When the cooling rate is less than 1 ° C./hour, the time required for cooling increases and the productivity decreases, so the cooling rate is set to 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 is transformed into pearlite, the hardness of the steel sheet is increased, and cold forgeability is deteriorated and impact resistance characteristics after carburizing and quenching and tempering are reduced. The cooling rate is 100 ° C./hour or less. Preferably, it is 90 ° C./hour.

  Here, the cooling stop temperature is a temperature to be controlled at the above cooling rate. If cooling to 650 ° C. is performed at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less, the temperature is reduced to 650 ° C. or less. The cooling is not particularly limited.

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

  As explained above, according to the manufacturing method that consistently manages the hot rolling conditions and annealing conditions of the present invention and performs the structure control of the steel sheet, it is excellent in cold forging combined with drawing and thickening forming. It is possible to produce a low carbon steel sheet that exhibits forgeability and also has excellent impact resistance after carburizing, quenching and tempering.

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

  A continuous cast slab (steel ingot) having the composition shown in Table 1 was heated at 1240 ° C. for 1.8 hours and then subjected to hot rolling. Finished hot rolling was completed at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C./second on the ROT, wound up at 510 ° C., and a hot-rolled coil having a thickness of 5.2 mm was manufactured.

  The hot-rolled coil is pickled, charged in a box-type annealing furnace, the atmosphere is controlled to 95% hydrogen-5% nitrogen, and then heated from room temperature to 705 ° C. at a heating rate of 100 ° C./hour. The temperature distribution in the coil was made uniform by maintaining at 705 ° C. for 36 hours. Then, after heating to 760 ° C. at a heating rate of 5 ° C./hour and further holding at 760 ° C. for 10 hours, cooling to 650 ° C. at a cooling rate of 10 ° C./hour, followed by furnace cooling to room temperature. A sample for characteristic evaluation was prepared.

  The sample structure was observed by the method described above, and the crack length existing in the sample after cold forging was measured by the method described above.

  Carburizing of the thickened sample was performed by gas carburization. In order to diffuse carbon from the atmosphere gas in the furnace to the inside of the steel through the sample surface layer, in the furnace in which the carbon potential is controlled to 0.5% by mass C, a treatment is performed at 940 ° C. for 120 minutes, It was cooled to the furnace.

  Subsequently, after heating from room temperature to 840 ° C., holding for 20 minutes was performed and quenching was performed in 60 ° C. oil. The quenched sample was subjected to a tempering treatment that was air-cooled after being held at 170 ° C. for 60 minutes to prepare a carburized quenched and tempered sample.

  The impact resistance of the carburized and tempered sample was evaluated by a drop weight test. FIG. 2 schematically shows an outline of a drop weight test for evaluating the impact resistance characteristics of a sample subjected to carburizing, quenching and tempering. The bottom of the cup-shaped sample 4 subjected to carburizing, quenching and tempering is fixed with a jig, and a weight of 2 kg is dropped on the cup side (upper side width: 50 mm, lower side width: 10 mm, height: 80 mm, length: 110 mm). Is dropped from the upper part 4 m away, an impact of about 80 J is applied to the vertical wall of the sample 4, the presence or absence of cracking of the sample is observed, and the impact resistance characteristics are evaluated.

  As a result of free fall, samples with no cracks or breakage were given an “OK” rating with excellent impact resistance, and samples with cracks or breakage were “NG” with poor impact resistance. Scored.

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

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

  On the other hand, the comparative steel L-1 has a low C amount and a hardness before cold forging of less than 100 HV, so that the cold forgeability is low. Comparative steels M-1, P-1, and Z-1 contain excessive amounts of P, Al, and N, and have a large amount of segregation at the γ / α interface during the second stage annealing. The formation of carbide is suppressed.

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

  Since the comparative steel O-1 has a small amount of Cr and the austenite grains on the surface layer are abnormally coarsened during carburizing, the impact resistance is low. Since comparative steel R-1 contains S excessively, coarse MnS produces | generates in steel and cold forgeability is low. Since the comparative steel U-1 contains C excessively, coarse carbides are generated in the thickened thickness of the steel and coarse carbides remain even after carburizing and quenching, so that the impact resistance is low.

  Since comparative steel V-1 had a small amount of Mn and it was difficult to increase the stability of carbide, cold forgeability and impact resistance after carburizing and tempering were low. Since the comparative steels W-1 and X-1 contain excessive amounts of O and Ti, the oxides and TiC present in the ferrite grains become carbide generation sites in the slow cooling after the two-phase annealing, and the grain boundaries The formation of carbides in is suppressed, and the cold forgeability is low. Since comparative steel Y-1 contains B excessively, cold forgeability is low.

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

  Table 4 shows the carbide diameter, pearlite area ratio, ferrite particle diameter, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the thickness of the vertical wall portion of the prepared sample. The ratio of the maximum crack length and the measurement results and evaluation results of impact resistance are shown.

  The comparative steel E-3 has a low finishing hot rolling temperature, an increased rolling load, and low productivity. Comparative Steel D-2 has a high finish hot rolling temperature and scale flaws formed on the surface of the steel sheet. Therefore, when subjected to a wear resistance test after quenching and tempering, cracks and peeling occurred starting from the scale flaws, resulting in wear resistance. The characteristics deteriorated. The comparative steel F-2 has a slow cooling rate in ROT (Run Out Table), which causes a decrease in productivity and generation of scale defects.

  In Comparative Steel C-4, the cooling rate at the ROT was 100 ° C./sec, and the outermost layer portion of the steel sheet was excessively cooled, so that fine cracks were generated in the outermost layer portion. The comparative steel C-2 has a low coiling temperature, a lot of low-temperature transformation structures such as bainite and martensite are generated and become brittle, cracks occur frequently when the hot-rolled coil is discharged, and productivity is lowered. Furthermore, the wear resistance characteristics of samples taken from the cracks are low.

  The comparative steel G-2 has a high coiling temperature, and a pearlite with a thick lamellar spacing is generated in the hot rolled structure, and the thermal stability of the needle-like coarse carbide is high, and after the two-step type annealing. However, since the carbides remain in the steel sheet, the machinability is low. Since the comparative steel H-4 has a slow heating rate in the first stage annealing of the two-step type annealing, the productivity is low.

  Since comparative steel E-3 has a high heating rate in the first stage annealing, the temperature difference between the inside of the coil and the inner and outer peripheral portions becomes large, and the soot and seizure due to the difference in thermal expansion occur, and quenching occurs. When subjected to an evaluation test of wear resistance after tempering, cracks and peeling occurred from the heel part, and the wear resistance was lowered.

  Comparative steel G-4 has a low holding temperature (annealing temperature) in the first stage annealing, the carbide coarsening treatment at the Ac1 point or less, and the thermal stability of the carbide is insufficient. Thus, the remaining carbides in the second stage annealing are reduced, and the pearlite transformation cannot be suppressed in the structure after the slow cooling, so that the machinability is low.

  Comparative steel D-4 has a high holding temperature (annealing temperature) in the first stage annealing, austenite is generated during annealing, and the stability of the carbide cannot be increased. Therefore, pearlite is generated after annealing, and Vickers hardness Is over 180 HV, and machinability is low. Comparative steel J-4 has a short holding time in the first-stage annealing, cannot increase the stability of the carbide, and has low machinability.

  Comparative steel F-2 has a long holding time in the first-stage annealing, low productivity, seizure flaws, and low wear resistance. Since the comparative steel B-4 has a low heating rate in the second-stage annealing of the two-step type annealing, the productivity is low. Since the comparative steel A-3 has a high heating rate in the second stage annealing, the temperature difference between the inside and the outer periphery of the coil is large, and the soot and seizure due to a large difference in thermal expansion due to transformation occurs. Low wear resistance after quenching and tempering.

  The comparative steel K-2 has a low holding temperature (annealing temperature) in the second stage annealing, a small amount of austenite generation, and cannot increase the number ratio of carbides at the ferrite grain boundaries, and therefore has low machinability. Since the comparative steel C-4 has a high holding temperature (annealing temperature) in the second stage annealing, and the dissolution of carbides during the annealing is promoted, it becomes difficult to form grain boundary carbides after slow cooling. Produced, Vickers hardness exceeds 180HV, and machinability is low.

  Since the comparative steel J-3 has a long holding time in the second stage annealing and promotes dissolution of carbides, the machinability is low. Comparative steel D-3 has a slow cooling rate from the second-stage annealing to 650 ° C., has low productivity, and coarse carbides are generated in the structure after slow cooling. Cracks occurred as a starting point, and cold forgeability deteriorated. Since the comparative steel I-3 has a high cooling rate from the second stage annealing to 650 ° C. and a pearlite transformation occurs during cooling to increase the hardness, the cold forgeability is low.

  Next, in order to investigate the allowable content of other elements, a continuous cast slab (steel ingot) having the composition shown in Table 5 and Table 6 (continuation of Table 5) was heated at 1240 ° C. for 1.8 hours, It was subjected to hot rolling. Finished hot rolling was completed at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C./second on the ROT, wound up at 510 ° C., and a hot-rolled coil having a thickness of 5.2 mm was manufactured.

  After pickling the hot-rolled coil, inserting the hot-rolled coil in a box-type annealing furnace, and controlling the atmosphere to 95% hydrogen-5% nitrogen, the temperature from room temperature to 705 ° C. was increased at a heating rate of 100 ° C./hour. Heat and hold at 705 ° C. for 36 hours to homogenize the temperature distribution in the coil, then heat to 760 ° C. at a heating rate of 5 ° C./hour, further hold at 760 ° C. for 10 hours, and then to 650 ° C. Was cooled at a cooling rate of 10 ° C./hour and then cooled to room temperature to prepare a sample for characteristic evaluation.

  In addition, the structure of the sample was observed by the above-described method, and the crack length existing in the sample after cold forging was measured by the above-described method.

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

  As shown in Table 7, invention steels AA-1, AB-1, AC-1, AD-1, AE-1, AF-1, AG-1, AH-1, AI-1, AJ-1, AK -1, AL-1, AM-1, AN-1, AO-1, AP-1, and AQ-1 are all the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains Is more than 1, Vickers hardness is 100HV or more and 180HV or less, and is excellent in cold forgeability and impact resistance after carburizing and tempering.

  On the other hand, comparative steels AR-1, AS-1, AW-1, AZ-1, BB-1, and BF-1 each contain excessive amounts of La, As, Cu, Ni, Sb, and Ce. In addition, since the amount of segregation at the γ / α interface increases during the second stage annealing, the formation of carbides at the grain boundaries is suppressed. Since comparative steel BG-1 contains Si excessively and the ductility of ferrite is low, cold forgeability is low.

  Since the comparative steels AT-1, AV-1, BA-1, BC-1, BH-1, and BJ-1 contain excessive amounts of Mo, Nb, Cr, Ta, W, and V, respectively, ferrite The carbide is finely dispersed in the grains, and the hardness exceeds 180 HV. Since comparative steel BF-1 contains Mn excessively, the impact resistance after carburizing and tempering is remarkably low.

  Comparative steels AU-1, AX-1, AY-1, and BE-1 contain excessive amounts of Zr, Ca, Mg, and Y, respectively, and coarse oxides or nonmetallic inclusions are produced in the steel. Then, cracks occurred starting from coarse oxides or coarse non-metallic inclusions during cold forging, and cold forgeability deteriorated. Comparative steel BD-1 contains excessive Sn, the ferrite becomes brittle, and the cold forgeability is low. Since the comparative steel BK-1 contains excessive C, coarse carbides are generated inside the thickened steel, and the coarse carbides remain even after carburizing and quenching, resulting in reduced impact resistance.

  Then, in order to investigate the influence of manufacturing conditions, AA, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP, and AQ shown in Table 5 A hot-rolled sheet annealing sample having a sheet thickness of 5.2 mm was produced from the slab having the composition of the following conditions under the hot-rolling conditions and annealing conditions shown in Table 8.

  Table 9 shows the carbide diameter, pearlite area ratio, ferrite particle diameter, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the maximum thickness of the vertical wall portion in the prepared sample. The ratio of crack length and the measurement results and evaluation results of impact resistance are shown.

  Comparative steel AC-2 has a low finishing hot rolling temperature and low productivity. The comparative steel AN-4 has a high finish hot rolling temperature, scales are formed on the steel sheet surface, and when an impact load is applied after cold forging and carburizing and quenching and tempering, cracks are generated from the flanges, resulting in impact resistance. The characteristics deteriorated.

  Inventive steel AB-3 has a slow cooling rate in the ROT, resulting in a decrease in productivity and the derivation of scale defects. Inventive steels AJ-3 and AD-4 had a cooling rate at ROT of 100 ° C./second, the outermost layer portion of the steel sheet was excessively cooled, and fine cracks were generated in the outermost layer portion.

  The comparative steel AN-3 has a low coiling temperature, a large amount of low-temperature transformation structures such as bainite and martensite are generated and become brittle, and cracks frequently occur when the hot-rolled coil is discharged, resulting in a reduction in productivity. Furthermore, the impact resistance characteristics after cold forging and carburizing quenching and tempering in samples taken from the cracked pieces were inferior.

  The comparative steel AH-3 has a high coiling temperature, and pearlite with thick lamellar spacing is generated in the hot rolled structure, and the thermal stability of the needle-like coarse carbide is high, even after annealing of the two-step type. Since the carbide remains in the steel sheet, the cold forgeability is low.

  Since the comparative steel AF-4 has a slow heating rate in the first stage annealing of the two-step type annealing, the productivity is low. Since the comparative steel AG-2 has a high heating rate in the first-stage annealing, the temperature difference between the inside and the outer periphery of the coil is large, so that the flaws and seizure due to the difference in thermal expansion occur, Impact resistance after forging and carburizing quenching and tempering decreased.

  The comparative steel AA-2 has a low holding temperature (annealing temperature) in the first stage annealing, the carbide coarsening treatment at the Ac1 point or less is insufficient, and the thermal stability of the carbide is insufficient. The carbide remaining at the time of annealing of the eyes decreased, the pearlite transformation could not be suppressed in the structure after the slow cooling, and the cold forgeability was lowered.

  The comparative steel AM-3 has a high holding temperature (annealing temperature) in the first stage, austenite is generated during annealing, and the stability of the carbide cannot be improved, and the cold forgeability and the resistance after carburizing and quenching and tempering are reduced. Impact characteristics deteriorated. Comparative steel AF-2 has a short holding time in the first-stage annealing, cannot improve the stability of carbide, and has low cold forgeability. Comparative steel AO-4 has a long holding time in the first stage annealing and low productivity.

  Since the comparative steel AP-4 has a slow heating rate in the second stage annealing of the two-step type annealing, the productivity is low. Since the comparative steel AI-3 has a high heating rate in the second stage annealing, the temperature difference between the inside and the outer periphery of the coil becomes large, and the soot and seizure due to a large difference in thermal expansion due to transformation occurs, and carburizing. When an impact load was applied after quenching and tempering, cracks were generated from the ridges, and the impact resistance characteristics deteriorated.

  The comparative steel AL-3 has a low holding temperature (annealing temperature) in the second stage annealing, a small amount of austenite is produced, the number ratio of carbides at the ferrite grain boundaries cannot be increased, and the cold forgeability is low. Declined. Since the comparative steel AD-2 has a high holding temperature (annealing temperature) in the second stage annealing and the dissolution of carbides is accelerated during annealing, it becomes difficult to generate grain boundary carbides after slow cooling, and cold forgeability. And the impact resistance after carburizing quenching and tempering decreased.

  Since the comparative steel AJ-4 has a long holding time in the second stage annealing and promotes dissolution of carbides, the cold forgeability is low. In comparison steel AQ-3, the cooling rate from the second stage annealing to 650 ° C. is slow, the productivity is low, and coarse carbides are generated in the structure after the slow cooling. As a starting point, cracks occurred and cold forgeability deteriorated. Since the comparative steel AP-2 has a high cooling rate from the second stage annealing to 650 ° C., pearlite transformation occurs during cooling, and the hardness increases, the cold forgeability deteriorates.

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

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

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

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

  As described above, according to the present invention, it is possible to provide a low carbon steel sheet excellent in cold forgeability and impact resistance after carburizing and tempering, and a method for producing the same. Since the steel sheet of the present invention is suitable as a material for obtaining parts such as high cycle gears by forming by cold forging such as sheet forming, the present invention has high industrial applicability. is there.

1 Disc-shaped test material 2 Cup-shaped test material 3 Crack 4 Sample 5 Falling weight L Maximum length of crack

(1) Component composition is mass%, C: 0.10-0.40%, Si: 0.01-0.30%, Mn: 0.30-1.00%, Al: 0.001- 0.10%, Cr: 0.50-2.00%, Mo: 0.001-1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020% or less O: 0.020% or less, Ti: 0.010% or less, B: 0.0005% or less, Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or less, Nb : 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0 0.050% or less, Ca: 0.050% or less, Y: 0.050% or less, Zr: 0.050% or less, La: 0.050% or less, and Ce: 0.05% or less. Comprises 50% or less, a balance being Fe and impurities der Ru steel plate,
Metal structure of the upper Symbol steel plate, pearlite area ratio is 6% or less, and the balance of ferrite and carbide, carbide grain size 0.4 to 2.0 .mu.m, the particle size of the ferrite is 3.0～50.0Myuemu, and the ferrite grain boundaries of the ratio greater than 1 of the number of carbides on the number of carbide in ferrite grains, the filled, Vickers hardness of the upper Symbol steel plate excellent in impact resistance, characterized in that at most 180HV than 100HV steel sheet.

(2) A production method for producing a steel plate having excellent impact resistance as described in (1) above, and finishing hot rolling of a steel slab having the component composition of (1) in a temperature range from 650 ° C. to 950 ° C. The hot-rolled steel sheet is subjected to hot rolling to obtain a hot-rolled steel sheet, the hot-rolled steel sheet is wound at 400 ° C. to 600 ° C., the hot-rolled steel sheet wound is pickled, and the pickled hot-rolled steel sheet is 30 ° C./hour. Heating to an annealing temperature of 650 ° C. or more and 720 ° C. or less at a heating rate of 150 ° C./hour or less and performing first-stage annealing for holding for 3 hours or more and 60 hours or less, Heating to an annealing temperature of 725 ° C. or more and 790 ° C. or less at a heating rate of 80 ° C./hour or more for a period of time, and performing a second stage annealing for holding for 3 hours or more and 50 hours or less, Cool to 650 ° C at a cooling rate of 1 ° C / hour or more and 100 ° C / hour or less. Steel sheet excellent manufacturing method of the cold forgeability and impact resistance characterized.

Claims (2)

Ingredient composition is mass%,
C: 0.10 to 0.40%,
Si: 0.01-0.30%,
Mn: 0.30 to 1.00%
Al: 0.001 to 0.10%,
Cr: 0.50 to 2.00%
Mo: 0.001 to 1.00%,
P: 0.020% or less,
S: 0.010% or less,
N: 0.020% or less,
O: 0.020% or less,
Ti: 0.010% or less,
B: 0.0005% or less,
Sn: 0.050% or less,
Sb: 0.050% or less,
As: 0.050% or less,
Nb: 0.10% or less,
V: 0.10% or less,
Cu: 0.10% or less,
W: 0.10% or less,
Ta: 0.10% or less,
Ni: 0.10% or less,
Mg: 0.050% or less,
Ca: 0.050% or less,
Y: 0.050% or less,
Zr: 0.050% or less,
La: 0.050% or less, and Ce: 0.050%
A low carbon steel plate that is the balance Fe and impurities,
The metal structure of the low carbon steel sheet is
Carbide particle size is 0.4-2.0 μm,
The pearlite area ratio is 6% or less, and the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains is more than 1,
The filling,
A steel sheet, wherein the low carbon steel sheet has a Vickers hardness of 100 HV or more and 180 HV or less. A manufacturing method for manufacturing the steel sheet according to claim 1,
The steel slab having the component composition according to claim 1 is subjected to hot rolling to complete hot rolling in a temperature range of 650 ° C. or higher and 950 ° C. or lower to obtain a hot rolled steel plate,
Winding the hot-rolled steel sheet at 400 ° C. or more and 600 ° C. or less,
The wound hot-rolled steel sheet is pickled, and the pickled hot-rolled steel sheet is heated to an annealing temperature of 650 ° C. to 720 ° C. at a heating rate of 30 ° C./hour to 150 ° C./hour for 3 hours. Apply the first stage of annealing for 60 hours or less, and then
The hot-rolled steel sheet is heated to an annealing temperature of 725 ° C. or more and 790 ° C. or less at a heating rate of 1 ° C./hour or more and 80 ° C./hour or less, and then subjected to second-stage annealing for 3 hours or more and 50 hours or less and annealing. A method for producing a steel sheet, comprising cooling a subsequent hot-rolled steel sheet to 650 ° C. at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less.

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