[0001]The present invention relates to a steel sheet and a manufacturing method thereof.
Background technique
[0002]By mass%, the steel sheet 0.1 to 0.4% carbon, from blank, press molding, hole expansion molding, bending, drawing, thickening and thinning the molding, and combinations thereof cold forging is performed, automotive gear, is used as the material of the drive system components such as a clutch. Compared to conventional hot forging and the like, the higher the amount of strain accumulated in the material by cold forging, lead to the occurrence of buckling when the material cracks and molding, there is a problem of causing deterioration of the component characteristics.
[0003]
　In particular, after performing carburizing hardening and tempering in the molding material in order to obtain the abrasion resistance, the residual stress by heat treatment is born in a state leading to generation and development of cracks from the crack portion and the buckling portion. For use as a drive system components, such as by biting Startup gear for momentarily large load bearing, since the acquisition of impact resistance for not brittle destruction is sought, the the steel sheet is required to secure the impact resistance after carburizing quenching and tempering and excellent cold forgeability.
[0004]
　Previously, many proposals for technology for improving the impact resistance after carburizing and cold forgeability of the steel sheet have been made (for example, Patent Documents 1-5, reference).
[0005]
　For example, Patent Document 1, as a machine structural steel with improved toughness by inhibiting the grain coarsening in carburizing heat treatment, by mass%, C: 0.10 ~ 0.30%, Si: 0.05 ~ 2.0%, Mn: 0.10 ~ 0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 ~ 3.00%, Al: 0.005 ~ 0.050%, Nb: 0.02 ~ 0.10%, N: containing 0.0300% or less, and the balance Fe and unavoidable impurities, cold working before the tissue is ferrite-pearlite structure, the mechanical structural steel is disclosed average of the ferrite grain diameter is 15μm or more.
[0006]
　Patent Document 2, as an excellent steel carburizing hardenability and cold workability, C: 0.15 ~ 0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, B: contains from 0.005 to 0.050%, the balance being Fe and unavoidable impurities, and the main ferrite phase and the graphite phase steel having a tissue is disclosed.
[0007]
　Patent Document 3, the steel for carburizing bevel gear having excellent impact strength, high toughness carburizing bevel gear, and its manufacturing method are disclosed.
[0008]
　Patent Document 4, after spheroidizing annealing, carried out cold forging, to parts produced by carburizing and quenching and tempering step, while having excellent processability, suppressing coarsening of crystal grains in the subsequent carburizing and excellent impact resistance, steel for carburized parts having impact fatigue characteristics is disclosed.
[0009]
　Patent Document 5, as plasma carburizing cold work tool steel, C: 0.40 ~ 0.80%, Si: 0.05 ~ 1.50%, Mn: 0.05 ~ 1.50%, and, V: 1.8 contained ~ 6.0%, further, Ni: 0.10 ~ 2.50%, Cr: 0.1 ~ 2.0%, and, Mo: 3.0% or less of one or comprise two or more, balance steel being Fe and inevitable impurities is disclosed.
CITATION
Patent Literature
[0010]
Patent Document 1: JP 2013-040376 Patent Publication
Patent Document 2: JP-A 06-116679 JP-
Patent Document 3: JP-A-09-201644 Publication
Patent Document 4: JP 2006-213951 Patent Publication
Patent Document 5: JP 10-158780 JP
Summary of the Invention
Problems that the Invention is to Solve
[0011]
　Organization of steel for mechanical structure of Patent Document 1 is a structure of ferrite + pearlite, the tissue, compared to a ferrite + cementite structure, because it has a large hardness, suppressing the wear of the mold in the cold forging can not, not necessarily the machine structural steel excellent in cold forgeability.
[0012]
　In the steel of Patent Document 2, the graphitization process of cementite annealing at a high temperature becomes necessary, it is impossible to suppress an increase in reduction and manufacturing cost of the yield.
[0013]
　Manufacturing method of Patent Document 3, it is necessary to perform a further hot forging after physical cold forging and carburizing, since hot forging is essential, not the production process leading to drastic cost reduction.
[0014]
　Carburized parts steel of Patent Document 4 is whether it is possible to achieve the same effect unclear in cold forging with a large strain is given, furthermore, a control method for specific organizational forms and organizations not clear is because, even in given forging molding a large strain in the cold, such as recently applied spreads plate forging can not be said steel exhibits excellent workability.
[0015]
　Patent Document 5, the formability of the steel, knowledge and techniques regarding optimal components and morphology of particularly to improve the cold forgeability is not any disclosure.
[0016]
　The present invention has been made in view of the circumstances of the prior art, excellent in impact resistance after carburizing quenching and tempering the cold forgeability, suitable steel and a manufacturing method thereof for obtaining a component such as high cycle gear in particular plate-shaped it is an object to provide.
Means for Solving the Problems
[0017]
　Wherein the foregoing problems, in order to obtain a steel sheet suitable for materials such as the drive system components, in a steel sheet containing the necessary C to increase the hardenability, increase the grain size of the ferrite, carbide (mainly cementite ) was spheronized in a suitable grain size, it can be understood that there may be fewer pearlite structure. This is due to the following reasons.
[0018]
　Ferrite phase has a low hardness, high ductility. Accordingly, ferrite mainly as tissue was, by increasing the grain size, it is possible to improve the material formability.
[0019]
　Carbides, by properly dispersed in the metallic structure, while maintaining the material formability, since it is possible to impart excellent wear resistance and rolling fatigue characteristics, must be in the drive system components tissue it is. Also, the carbide in the steel sheet is a strong particles that prevent slipping, the presence of the carbides in the ferrite grain boundaries, to prevent the propagation of slip exceeding a crystal grain boundary, to suppress the formation of shear bands can, to improve the cold forgeability, at the same time, the moldability of the steel sheet improved.
[0020]
　However, cementite is hard and brittle structure, when present in pearlite state a layered structure of ferrite, steel hard, so becomes brittle, should be present in a spherical shape. And cold forgeability, considering the occurrence of cracks during forging, the particle size should be appropriate range.
[0021]
　However, the manufacturing method for realizing the tissue has not been previously disclosed. Accordingly, the present inventors have conducted extensive studies on manufacturing method for realizing the above tissues.
[0022]
　As a result, since the cementite metal structure of the steel sheet after coiling after hot rolling in small fine pearlite or fine ferrite having lamella spacing is the dispersed bainite, a relatively low temperature (400 ℃ ~ 550 ℃) in the take-take. By winding at a relatively low temperature, cementite dispersed in ferrite also easily spheroidized. Subsequently, partially spheroidized cementite with annealing at temperatures just below Ac1 point as annealing in the first stage. Then, in annealing at a temperature between Ac1 point and Ac3 points as annealing in the second stage (two-phase region of the so-called ferrite and austenite), while leaving a portion of the ferrite grains, thereby austenite transformation part. While growing the subsequent slow cooling to leave ferrite grains, by which it to ferrite transformation of austenite in the nucleus, grain boundary precipitation of cementite while obtaining large ferrite phase, it was found to be able to realize the tissue.
[0023]
　That is, the manufacturing method of the steel sheet to satisfy the hardenability formability at the same time, such as hot rolling conditions and annealing conditions are also difficult to implement to devise in a single, in the so-called integrated process, such as hot rolling, annealing step It was found that can be realized by achieving an optimization.
[0024]
　Further, for improving the drawability of the cold forging is necessary to reduce the plastic anisotropy, this improvement, the adjustment of the hot rolling condition is found that is important.
[0025]
　The present invention has been accomplished based on these findings, the gist is as follows.
[0026]
　(1) component composition, by 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 .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. Comprises 50%, a low carbon steel sheet which is a balance of Fe and impurities, metal structure of the low-carbon steel sheet, carbide particle size of 0.4 ~ 2.0 .mu.m, pearlite area ratio is 6% or less, and ferrite grains steel the ratio of the number of ferrite grain boundary carbide to the number of the carbide of the inner satisfies 1 greater than a, and wherein the Vickers hardness of the low carbon steel is less than 180HV than 100 HV.
[0027]
　(2) the A method of manufacturing a steel plate (1), wherein (1) the billet chemical composition subjected to complete hot rolling finishing hot-rolling in a temperature range of 650 ° C. or higher 950 ° C. or less of Te and hot-rolled steel sheet, wound up the hot-rolled steel sheet at 400 ° C. or higher 600 ° C. or less, wound hot rolled steel sheets were subjected to pickling, the pickled hot-rolled steel sheet 30 ° C. / time or 0.99 ° C. / time or less in the heating rate, 650 ° C. by heating above 720 ° C. below the annealing temperature, annealed at the first stage to 3 hours or more 60 hours or less, then, 80 ° C. / time 1 ° C. / time or a hot-rolled steel sheet the following heating rate, heated to an annealing temperature of 790 ° C. or less 725 ° C. or higher, annealed in the second stage to hold 3 hours or more 50 hours or less, the hot-rolled steel sheet after annealing, 1 ° C. / time more than 100 cold, characterized in that cooling to 650 ° C. at ° C. / time less cooling rate Manufacturing method of forging property and steel plate.
Effect of the invention
[0028]
　According to the present invention, excellent impact properties after cold forgeability and carburizing quenching and tempering, it is possible to provide a suitable steel and its manufacturing method to obtain particularly parts such as high cycle gear by sheet forming.
BRIEF DESCRIPTION OF THE DRAWINGS
[0029]
FIG. 1 is a diagram schematically showing the manner of the introduced crack Overview and cold forging cold forging test. (A) shows a disk-shaped test piece cut from hot rolled steel sheet, (b) shows the shape of the test material after cold forging, (c) is a cross-sectional aspect of the test material after cold forging It is shown.
2 is a diagram schematically showing an outline of drop weight test for evaluating the impact resistance of the samples subjected to carburizing quenching and tempering.
3 is a diagram showing the ratio of the number of grain boundary carbides to the number of grain carbide, the relationship between the crack length and the impact resistance after carburizing quenching and tempering of the cold forging test piece.
4 is a diagram showing the ratio of the number of grain boundary carbides to the number of grain carbide, another relationship crack length and impact resistance after carburizing quenching and tempering of the cold forging test piece.
DESCRIPTION OF THE INVENTION
[0030]
　Hereinafter, the present invention will be described in detail. First, a description will be given reasons for limiting the chemical composition of the steel sheet of the present invention. Here, according to the chemical composition "%" means "% by mass".
[0031]
　[C: 0.10 ~ 0.40%]
　C is a carbide is formed in the steel is an element effective for refining the strengthening and ferrite grains of the steel. Suppressing the occurrence of satin in the cold working, to ensure the surface appearance of the cold forging is the suppression of coarsening of the ferrite grain size is essential, it is less than 0.10%, the volume percentage of carbide There was insufficient, since it is impossible to suppress the coarsening of carbides during annealing, C is 0.10% or more. Preferably 0.11% or more.
[0032]
　On the other hand, when it exceeds 0.40%, the increase in the volume percentage of carbide, momentarily crack as a starting point of fracture at the time of a load is generated in a large amount, so deteriorating the impact resistance, C is 0.40% or less. Preferably not more than 0.38%.
[0033]
　[Si: 0.01 ~
　0.30%] Si acts as a deoxidizer and is also affects elements in the form of carbides. Reducing the number of carbide in ferrite grains to obtain a deoxidation effect, in order to increase the number of carbide on the grain boundary of ferrite is the annealing of the two-stage stepped, during annealing, to generate austenite phase, once, gradually cooled after dissolving the carbides, it is necessary to promote carbide formation in the ferrite grain boundaries.
[0034]
　When Si exceeds 0.30%, the ferrite ductility is lowered, is likely to occur cracking during cold forging, because the impact resistance after carburizing quenching and tempering the cold forgeability is degraded, Si 0.30 % or less to be. Preferably not more than 0.28%.
[0035]
　Si is preferably as small, reduced to less than 0.01%, since leads to a significant increase in refining costs, Si is 0.01% or more. Is preferably 0.02% or more.
[0036]
　[Mn: 0.30 ~
　1.00%] Mn, in the annealing of the two-stage stepped, an element for controlling the form of carbides. If it is less than 0.30%, the slow cooling after the second stage annealing, since it is difficult to produce a carbide on the grain boundary of ferrite, Mn is set to 0.30% or more. Preferably 0.33% or more.
[0037]
　On the other hand, when it exceeds 1.00%, the toughness after carburizing quenching and tempering is deteriorated, Mn is not more than 1.00%. Preferably not more than 0.96%.
[0038]
　[Al: 0.001 ~
　0.10%] Al is an element to stabilize the ferrite acts as a deoxidizer for steel. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Al is 0.001% or more. Preferably 0.004% or more.
[0039]
　On the other hand, when it exceeds 0.10%, reduces the number ratio of carbides on the grain boundaries, so causing an increase in crack length during cold forging, Al is not more than 0.10%. Preferably not more than 0.09%.
[0040]
　[Cr: 0.50 ~
　2.00%] Cr and Mo is an element that improves the toughness. Cr is an effective element in the stabilization of carbides during the heat treatment. If it is less than 0.50%, it is difficult to leave the carbides during carburizing, leads to coarsening of austenite grain size in the surface layer, so causing a decrease in impact resistance, Cr is 0.50% or more. Preferably 0.52% or more.
[0041]
　On the other hand, when it exceeds 2.00%, an increase in the concentrated amount of Cr in the carbides, the austenite phase generated by annealing two-step type, for the remaining number fine carbide grain after annealing also present carbides within, reduces the number ratio of hardness increase and grain boundary carbides, since cold forgeability is degraded, Cr is not more than 2.00%. Preferably not more than 1.94%.
[0042]
　[Mo: 0.001 ~
　1.00%] Mo is an element effective for morphology control of carbides. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Mo is 0.001% or more. Preferably it is greater than or equal to 0.017%.
[0043]
　On the other hand, when it exceeds 1.00%, Mo is concentrated in carbides, for stable carbide even austenite phase increases after annealing exist carbides in the grains, the hardness increased and the grain boundary carbide cause a decrease in the number ratio, since cold forgeability is degraded, Mo is not more than 1.00%. Preferably not more than 0.94%.
[0044]
　The following elements are impurities, it should be controlled below a certain amount.
[0045]
　[P: 0.020% or less]
　P segregates in ferrite grain boundaries, a element for suppressing generation of grain boundary carbides. The less preferred. The content of P may be 0, but to highly purified to less than 0.0001% in refining process, a long time is required for refining, so causing a significant increase in manufacturing costs, a substantial lower limit it is 0.0001 to 0.0013%.
[0046]
　On the other hand, if it exceeds 0.020%, the decreased number ratio of the grain boundary carbides, since cold forgeability is degraded, P is the 0.020% or less. Preferably not more than 0.018%.
[0047]
　[S: 0.010% or less]
　S is an impurity element that forms non-metallic inclusions such as MnS. Non-metallic inclusions, since the starting point of cracking during cold forging, S as less preferred. The content of S may be 0, but when reducing the S to less than 0.0001%, the refining costs increase significantly, substantial lower limit is from 0.0001 to 0.0012%.
[0048]
　On the other hand, if it exceeds 0.010%, so cause an increase in crack length during cold forging, S is 0.010% or less. Preferably not more than 0.009%.
[0049]
　[N: 0.020% or less]
　N segregates to the grain boundary of ferrite is element for suppressing generation of carbides on grain boundaries. The less preferred. The content of N may be 0, but when reduced to less than 0.0001%, the refining costs increase significantly, substantial lower limit is 0.0001 to 0.0006%.
[0050]
　On the other hand, if it exceeds 0.020% and subjected to 2-phase region annealing and slow cooling, the ratio of the number of carbide on the grain boundary of ferrite to the number of carbide in ferrite grains is less than 1, the cold forgeability since lowered, N is the to 0.020% or less. Preferably is less than or equal to 0.017%.
[0051]
　[O: 0.0001 ~ 0.020%]
　O is an element which forms oxides in the steel. Oxide present in the ferrite grains is, since the carbide formation site, lesser is preferable. The content of O may be 0, but when to reduce the O to less than 0.0001%, the refining costs increase significantly, substantial lower limit is 0.0001 to 0.0006%.
[0052]
　On the other hand, if it exceeds 0.020% or the ratio of the number of carbide on the grain boundary of ferrite to the number of carbide in ferrite grains is less than 1, since the cold forgeability is degraded, O is the 0.020% or less to. Preferably is less than or equal to 0.017%.
[0053]
　[Ti: 0.010% or less]
　Ti is an important element to control the form of carbides, the content of a large amount, an element to promote the formation of carbides in the ferrite grain, the less preferred. The content of Ti may be 0, but when reduced to less than 0.0001%, the refining costs increase significantly, substantial lower limit is from 0.0001 to 0.0006%.
[0054]
　On the other hand, if it exceeds 0.010%, the ratio of the number of carbide on the grain boundary of ferrite to the number of carbide in ferrite grains is less than 1, since the cold forgeability is degraded, Ti is a 0.010% or less to. Preferably 0.007% or less.
[0055]
　[B: 0.0005% or less]
　B is an effective element for controlling a slip dislocation during cold forging. By containing a large amount of, the activities of the slip systems is limited, B is lesser preferred. The content of B may be 0. As well as a careful attention must be paid to the detection of less than 0.0001% B, by analysis apparatus, leading to below the detection limit.
[0056]
　On the other hand, if it exceeds 0.0005%, inhibited cross slip dislocations at a shear zone formed by cold forging, because cracks are concentrated locally distortion, B is not more than 0.0005% . Preferably not more than 0.0005%.
[0057]
　[Sn: 0.050% or
　less] Sn is an element which is mixed from the steel material (scrap), the less preferred. May be any 0 content of Sn, when reduced to less than 0.001%, the refining costs increase significantly, substantial lower limit is from 0.001 to 0.002%.
[0058]
　On the other hand, if it exceeds 0.050% ferrite is brittle, since the cold forgeability is degraded, Sn is set to 0.050% or less. Preferably is less than or equal to 0.048%.
[0059]
　[Sb: 0.050% or
　less] Sb, like Sn, an element which is mixed from the steel raw material (scraps). Sb is segregated at grain boundaries, as it reduces the number ratio of the grain boundary carbides, the less preferred. The content of Sb may be 0, but when reduced to less than 0.001%, the refining costs increase significantly, substantial lower limit is from 0.001 to 0.002%.
[0060]
　On the other hand, if it exceeds 0.050%, the cold forgeability is degraded, Sb is set to 0.050% or less. Preferably is less than or equal to 0.048%.
[0061]
　[As: 0.050% or
　less] As is, Sn, As with Sb, is an element which is mixed from the steel raw material (scraps), As is segregated at grain boundaries, as it reduces the number ratio of the grain boundary carbide , the less preferable. The content of As may be 0, but when reduced to less than 0.001%, the refining costs increase significantly, substantial lower limit is from 0.001 to 0.002%.
[0062]
　On the other hand, if it exceeds 0.050%, the decreased number ratio of the grain boundary carbides, since cold forgeability is degraded, As is the 0.050% or less. Preferably not more than 0.045%.
[0063]
　The present invention steel is a basic element of the above elements, further, cold forgeability and, in order to improve other properties, it may contain the following elements. The following elements, since it is not essential for obtaining the effects of the present invention, the content may be 0.
[0064]
　[Nb: 0.10% or less]
　Nb is an element effective for morphology control of carbides, also tissue is miniaturized, an element which contributes to the improvement of toughness. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Nb is preferably 0.001% or more. More preferably 0.002% or more.
[0065]
　On the other hand, when it exceeds 0.10%, precipitation many fine Nb carbide, strength is excessively increased, also reduces the number ratio of the grain boundary carbides, since cold forgeability is degraded, Nb is 0 and .10% or less. Preferably not more than 0.09%.
[0066]
　[V: 0.10% or less]
　V, similarly to the Nb, is an effective element to form the control of carbides, also tissue is miniaturized, an element which contributes to the improvement of toughness. If it is less than 0.001%, because the addition effect is not sufficiently obtained, V is preferably 0.001% or more. More preferably 0.004% or more.
[0067]
　On the other hand, when it exceeds 0.10%, precipitates fine V carbide are many, the strength is excessively increased, also reduces the number ratio of the grain boundary carbides, since cold forgeability is degraded, V 0 and .10% or less. Preferably not more than 0.09%.
[0068]
　[Cu: 0.10% or
　less] Cu forms fine precipitates, an element which contributes to improvement in strength. If it is less than 0.001%, the strength improvement effect is not sufficiently obtained, Cu is preferably 0.001% or more. More preferably 0.008% or more.
[0069]
　On the other hand, when it exceeds 0.10%, the expressed hot shortness during hot rolling, productivity Cu is not more than 0.10% since the decrease. Preferably not more than 0.09%.
[0070]
　[W: 0.10% or less]
　W also, Nb, similarly to V, is an element effective to form the control of the carbides. If it is less than 0.001%, because the addition effect is not sufficiently obtained, W is preferably set to 0.001% or more. More preferably 0.003% or more.
[0071]
　On the other hand, when it exceeds 0.10%, precipitation many fine W carbides, strength is excessively increased, also reduces the number ratio of the grain boundary carbides, since cold forgeability is degraded, W 0 and .10% or less. Preferably not more than 0.08%.
[0072]
　[Ta: 0.10% or
　less] Ta also, Nb, V, similarly to the W, is an element effective to form the control of the carbides. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Ta is preferably set to 0.001% or more. Preferably 0.007% or more.
[0073]
　On the other hand, when it exceeds 0.10%, precipitation many fine W carbides, strength is excessively increased, also reduces the number ratio of the grain boundary carbides, since cold forgeability is degraded, Ta is 0 and .10% or less. Preferably not more than 0.09%.
[0074]
　[Ni: 0.10% or
　less] Ni is an element effective in improving the impact resistance of the parts. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Ni is preferably 0.001% or more. More preferably 0.002% or more.
[0075]
　On the other hand, when it exceeds 0.10%, it decreases the number ratio of the grain boundary carbides, since cold forgeability is degraded, Ni is not more than 0.10%. Preferably not more than 0.09%.
[0076]
　[Mg: 0.050% or
　less] Mg is an element capable of controlling the form of sulfide by the addition of small amount. If less than 0.0001%, the addition effect is not sufficiently obtained, Mg is preferably 0.0001% or more. More preferably not less than 0.0008%.
[0077]
　On the other hand, if it exceeds 0.050% ferrite is brittle, since the cold forgeability is degraded, Mg is set to 0.050% or less. Preferably is less than or equal to 0.049%.
[0078]
　[Ca: 0.050% or
　less] Ca, like Mg, an element capable of controlling the form of sulfide by the addition of small amount. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Ca is preferably 0.001% or more. More preferably 0.003% or more.
[0079]
　On the other hand, if it exceeds 0.050% produces coarse Ca oxides, since the starting point of cracking during cold forging, Ca is set to 0.050% or less. Preferably not more than 0.04%.
[0080]
　[Y: 0.050% or less]
　Y is, Mg, as with Ca, an element that can control the form of sulfide by the addition of small amount. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Y is preferably 0.001% or more. More preferably 0.003% or more.
[0081]
　On the other hand, if it exceeds 0.050% produces coarse Y oxides, since the starting point of cracking during cold forging, Y is a 0.050% or less. Preferably is less than or equal to 0.031%.
[0082]
　[Zr: 0.050% or
　less] Zr is, Mg, Ca, similar to the Y, is an element capable of controlling the form of sulfide by the addition of small amount. If it is less than 0.001%, because the addition effect is not sufficiently obtained, Zr is preferably set to 0.001% or more. More preferably 0.004% or more.
[0083]
　On the other hand, if it exceeds 0.050% produces coarse Zr oxides, since the starting point of cracking during cold forging, Zr is set to 0.050% or less. Preferably not more than 0.045%.
[0084]
　[La: 0.050% or less]
　La is an effective element to form the control of sulfide by the addition of trace amounts, also segregated in the grain boundary, is an element to reduce the number ratio of grain boundary carbides. If it is less than 0.001%, since the shape control effect is not sufficiently obtained, La is preferably 0.001% or more. More preferably 0.003% or more.
[0085]
　On the other hand, if it exceeds 0.050%, the decreased number ratio of the grain boundary carbides, since cold forgeability is degraded, La is set to 0.050% or less. Preferably is less than or equal to 0.047%.
[0086]
　[Ce: 0.050% or less]
　Ce, like La, and an element capable of controlling the form of sulfide by the addition of trace amounts, also segregates at grain boundaries, reducing the number ratio of the grain boundary carbide element it is. If it is less than 0.001%, since the shape control effect is not sufficiently obtained, Ce is preferably 0.001% or more. More preferably 0.003% or more.
[0087]
　On the other hand, if it exceeds 0.050%, the decreased number ratio of the grain boundary carbides, since cold forgeability is degraded, Ce is the 0.050% or less. Preferably is less than or equal to 0.046%.
[0088]
　Incidentally, the balance of the chemical composition of the steel sheet of the present invention is Fe and inevitable impurities.
[0089]
　Next, a description will be given tissue of the steel sheet of the present invention.
[0090]
　Organization of the steel sheet of the present invention are substantially, tissue composed of ferrite and carbides. Carbides, cementite is a compound of iron and carbon (Fe 3 was added to C), the compound of Fe atoms in cementite Mn, was replaced with Cr or the like, alloy carbides (M 23 C 6 , M 6 C, be a MC or the like , M is Fe and other metal elements).
[0091]
　When forming the steel sheet into a predetermined product shape, shear zones are formed in the macrostructure of the steel sheet, in the vicinity of the shear zone, caused by slip deformation is concentrated. Slip deformation is accompanied by proliferation of dislocation, in the vicinity of the shear zone, the region of high dislocation density is formed. With the increase of the strain amount applied to the steel sheet, slip deformation is promoted, the dislocation density increases.
[0092]
　In cold forging, large deformation is applied to more than 1 in equivalent strain. Therefore, in the conventional steel sheet, it is impossible to prevent the occurrence of voids and / or cracks with the increase in dislocation density, it fades improved cold forgeability was difficult.
[0093]
　To solve this difficult problem, it is effective to suppress the formation of shear zones during molding. In terms of microstructure, the formation of shear bands, can be understood as a phenomenon of slippage generated in one certain grain continuously propagated to neighboring grains over the grain boundaries. Therefore, in order to suppress the formation of shear bands, it is necessary to prevent the propagation of slip exceeding a grain boundary.
[0094]
　Carbides in the steel sheet is a strong particles that prevent slipping, carbides, the presence in the ferrite grain boundaries to suppress the formation of shear bands, it is possible to improve the cold forgeability.
[0095]
　In order to obtain such an effect, carbides, it is necessary to disperse in an appropriate size in the metal structure. Therefore, the average particle diameter of carbide to 0.4μm or 2.0μm below. When the particle diameter of carbide is less than 0.4 .mu.m, the hardness of the steel sheet is increased significantly, the cold forgeability is degraded. More preferably 0.6μm or more.
[0096]
　On the other hand, when the average particle diameter of the carbide is more than 2.0 .mu.m, in the cold-forming, carbides it becomes a starting point of cracking. More preferably not more than 1.95μm.
[0097]
　Furthermore, cementite is carbide iron is hard and brittle structure, when present in pearlite state a layered structure of ferrite, hard and steel becomes brittle. Therefore, perlite must be minimized, in the steel sheet of the present invention, 6% or less in area ratio.
[0098]
　Perlite because of its unique lamellar structure, it is possible distinguish SEM, an optical microscope observation. By calculating the area of ​​the lamellar structure in any cross-section, it is possible to determine the area ratio of pearlite.
[0099]
　Based on theory and principles, cold forgeability is considered to strongly affected by the coverage of the ferrite grain boundary carbide, the the high-precision measurement is required, carbides into the ferrite grain boundaries in the three-dimensional space of the measurement of coverage, serial sectioning SEM observation performed with a scanning electron microscope in the repeated sample cutting and observation by FIB, or, 3D EBSP observation is essential, together requires enormous measurement time, the technical know-how accumulation is essential.
[0100]
　The present inventors have revealed that this, more simply the result of searching the high metric accuracy, if the ratio of the number of ferrite grain boundary carbide to the number of carbide in ferrite grains as an index, cold can be evaluated forgeability, the ratio of the number of the ferrite grain boundary carbide to the number of carbide in ferrite grains exceeds 1, that cold forging is remarkably improved, the present inventors have found .
[0101]
　Incidentally, buckling of the steel sheet caused during cold forging, folding, none of the convolution, since those caused by localized distortion due to the formation of the shear zone, likewise, be present carbides in the ferrite grain boundary in, when relaxing the localization of formation and distortion of the shear zone, it is possible to suppress buckling, folding, generation of convolution.
[0102]
　Observation of carbide is carried out with a scanning electron microscope. Prior to observation, the sample for microstructure observation was polished by diamond abrasive grains having an average particle size of wet grinding and 1μm by emery paper, after finishing the observation plane mirror, a saturated picric acid - the tissue at alcohol solution keep etching.
[0103]
　Observation magnification was 3000 times, randomly to eight photographs a view of 30 [mu] m × 40 [mu] m in the sheet thickness 1/4 layers. The obtained tissue image, the image analysis software represented by Mitani Corp. (Win ROOF), measuring the area of ​​each carbide contained in that region in detail. Determined circle equivalent diameter (= 2 × √ (area /3.14)) from the area of ​​the carbides, and the average value and the carbide particle size.
[0104]
　Incidentally, in order to suppress the influence of a measurement error due to noise, the area is 0.01 [mu] m 2 is less carbide excluded from the evaluation.
[0105]
　Counting the number of cementite present in the ferrite grain boundaries on, and subtracting the number of carbides on grain boundaries from all the carbides number, finding the number of carbide in ferrite grains. Based on the measured number, to calculate the number ratio of the grain boundary carbide against carbide in ferrite grains.
[0106]
　As structure after annealing, the ferrite grain size by a 3.0μm or 50.0μm or less, it is possible to improve the cold forgeability. When the ferrite grain diameter is less than 3 [mu] m, to increase the hardness, a crack or cracks in the cold forging is likely to occur, the ferrite grain size is more preferably 3.0 [mu] m. More preferably not less than 7.5μm.
[0107]
　On the other hand, when the ferrite grain diameter exceeds 50.0 micrometers, decreased number of carbides suppressing grain boundary propagation of slip, because the cold forgeability is degraded, ferrite grain size is preferably not more than 50.0 micrometers . More preferably less than or equal to 37.9μm.
[0108]
　Ferrite grain size is in the above procedure, after polishing the observation surface of the sample for microstructure observation mirror, 3% nitric acid - a viewing surface which is etched in alcoholic solution tissues were observed under an optical microscope or a scanning electron microscope , measured by applying a line segment method with respect to the photographed image.
[0109]
　The Vickers hardness of the steel sheet by a 100HV or more 180HV or less, it is possible to improve the impact resistance after cold forgeability and carburizing quenching and tempering. When Vickers hardness is less than 100HV, easily buckling occurs during cold forging, the convolution fold and buckling portion is lowered impact resistance occurs, Vickers hardness of 100HV or more and to. Preferably is equal to or greater than 110HV.
[0110]
　On the other hand, when the Vickers hardness exceeds 180 Hv, ductility is reduced, internal cracks is likely to occur during cold forging, because the impact resistance is deteriorated, Vickers hardness is not more than 180 Hv. Preferably it is less than 170HV.
[0111]
　Next, a description will be given of an evaluation method of cold forging properties.
[0112]
　Figure 1 shows an embodiment of a crack introduced in overview and cold forging cold forging test schematically. In FIG. 1 (a), shows a disk-shaped test piece cut from hot rolled steel sheet, in FIG. 1 (b), shows the shape of the test material after cold forging, in FIG. 1 (c), after cold forging It shows a cross-sectional aspect of the test material.
[0113]
　As shown in FIG. 1, the hot-rolled steel sheet having a thickness of 5.2 mm, it cuts out a disc-shaped test material 1 having a diameter of 70 mm (see FIG. 1 (a),), by deep drawing, the diameter of the bottom 30mm cupped the test material is prepared (not shown). Next, using the one-shot foaming press made forest Tekko, the vertical wall portion of the cup-shaped test material, thickened molding with the thickening ratio 1.54 (= 8 mm / 5.2 mm) to (cold forging) , diameter 30 mm, height 30 mm, to prepare a cup-shaped test material 2 TatekabeAtsu 8 mm (FIG. 1 (b), the reference).
[0114]
　The cup-shaped test material 2 subjected to thickness increasing molding at FANUC Ltd. wire cut electric discharge machine, cut to appear sectional diameter portions (FIG. 1 (c), the reference). The cut surface was mirror polished up to confirm that the presence crack 3, the ratio of the maximum length L of the crack that is present in the vertical wall portion to the thickness of the vertical wall portion after thickening (= crack the cut surface 8mm thick vertical wall portion after length L / thickening) is measured. This measure, evaluate the cold forgeability.
[0115]
　Even if the initial thickness is other than 5.2 mm, so that the height of the vertical wall after thickening is 30 mm, by adjusting the diameter of the disk-shaped test material to be cut out, with the thickening ratio of the same 1.54 if molded, regardless of the initial thickness, it is possible to reproduce the results of evaluation, the hot-rolled steel sheet covered by the present invention is not limited to the hot-rolled steel sheet having a thickness of 5.2 mm. The present invention also in the hot-rolled steel sheet of the general thickness (2 ~ 15 mm), it is possible to improve the impact resistance after carburizing quenching and tempering the cold forgeability.
[0116]
　Next, the present invention will be described manufacturing method. Technical idea of ​​the present invention production process, when producing a steel plate from a steel slab of the aforementioned component composition, to manage consistently hot rolling conditions and annealing conditions, the cold forgeability and impact resistance after carburizing quenching and tempering it is to improve.
[0117]
　Features of the present invention will be described manufacturing method.
[0118]
　Features of the hot rolled]
　Molten steel having the required chemical composition is continuously cast to a slab, the slab, conventional manner, as it is subjected to hot rolling, or once heated after cooling, subjected to hot rolling , and it ends the finish hot rolled in a temperature range of 650 ° C. or higher 950 ° C. or less. The hot-rolled steel sheet after finish rolling cooled on ROT, wound at 600 ° C. inclusive coiling temperature 400 ° C..
[0119]
　[Annealing Features]
　in the hot-rolled steel sheet, pickling, but subjected to annealing of 2-step type of holding at two temperature ranges, in which, 30 in the annealing of the first stage, the hot-rolled steel sheet, to the annealing temperature ° C. / heating time more than 0.99 ° C. / time less heating rate, subjected to holding annealing at a temperature range of 650 ° C. or higher 720 ° C. or less for 3 hours or more 60 hours or less.
[0120]
　In annealing the following two-stage, the hot-rolled steel sheet, annealing temperature to heating at 1 ° C. / time or 80 ° C. / time following a heating rate of 50 hours or more 3 hours to a temperature range of 790 ° C. or less 725 ° C. or higher subjected to annealing to hold below.
[0121]
　Then, the hot-rolled steel sheet after annealing, to 650 ° C., and cooled at 100 ° C. / time less than the cooling rate of 1 ° C. / time, then cooled to room temperature.
[0122]
　The cooperation of the hot rolling conditions and annealing conditions, it is possible to obtain a low-carbon steel sheet having excellent impact resistance after cold forgeability and carburizing quenching and tempering.
[0123]
　The following specifically describes the process conditions of the present invention production process.
[0124]
　[Hot rolling]
　　finishing hot-rolling temperature: 650 ° C. or higher 950 ° C. or less
　　coiling temperature: 400 ° C. or higher 600 ° C. or less
[0125]
　The molten steel having the required chemical composition is continuously cast to a slab, as it is or once heated after cooling, subjected to hot rolling, and ends the finish hot rolled in a temperature range of 650 ° C. or higher 950 ° C. or less, heat the rolled steel sheet wound at 400 ° C. or higher 600 ° C. or less.
[0126]
　Slab heating temperature is preferably 1300 ° C. or less, the heating time of the temperature of the slab surface is kept above 1000 ° C. is preferably below 7 hours.
[0127]
　The heating temperature exceeds 1300 ° C., or, if the heating time exceeds 7 hours, becomes significant decarburization slab surface, when before quenching heating, the surface layer of the austenite grains abnormally grow, lowered impact resistance since, the heating temperature is preferably 1300 ° C. or less, the heating time is preferably less than 7 hours. More preferably, the heating temperature is 1280 ° C. or less, the heating time is less than 6 hours.
[0128]
　Finishing hot rolling is terminated at a temperature of less than or equal to 950 ℃ 650 ℃ or more. When finishing hot-rolling temperature is lower than 650 ° C., from the increase in deformation resistance of the steel, increased remarkably the rolling load, further, since roll wear amount is increased, productivity is lowered, the finishing hot-rolling temperature 650 ° C. and more. Preferably at 680 ℃ or more.
[0129]
　On the other hand, when the finishing hot-rolling temperature exceeds 950 ° C., ROT thick scale is formed during passage through the (Run Out Table), due to the scale defects are generated on the surface of the steel sheet, the time of cold forging, and / or, when an impact load is applied after carburizing quenching and tempering, since the impact resistance decreases cracking is a flaw as a starting point, finishing hot rolling temperature is set to 950 ° C. or less. It is preferably at most 920 ° C..
[0130]
　Cooling rate during cooling of the hot-rolled steel sheet on the ROT is preferably 10 ° C. / sec or higher 100 ° C. / sec or less. When the cooling rate is less than 10 ° C. / sec, in the course cooling, and generation of thick scale, can not suppress the occurrence of scratches caused thereby, since impact resistance is lowered, the cooling rate is 10 ° C. / more seconds is preferred. More preferably 20 ° C. / sec or more.
[0131]
　On the other hand, over the inside from the surface layer of the steel sheet and cooling the hot-rolled steel sheet at a cooling rate exceeding 100 ° C. / sec, the outermost layer is excessively cooled, the outermost layer, the low-temperature transformation structure such as bainite or martensite occur.
[0132]
　After winding, when paying out a 100 ° C. ~ hot rolled steel sheet at room temperature, fine cracks are generated in the low temperature transformation structure, it is difficult to remove the cracks in the subsequent pickling and cold rolling step, the cold forging and / or when an impact load is applied after carburizing quenching and tempering, cracks cracks developed starting from the so deteriorating the impact resistance, the cooling rate is preferably 100 ° C. / sec or less. More preferably 80 ° C. / sec or less.
[0133]
　Incidentally, the cooling rate after the hot-rolled steel sheet after the finish hot rolling has passed the no-injection period, from the time of receiving water cooled in the water injection section, at the time it is cooled on ROT to the target temperature of the winding, each points to a cooling power received from the cooling facility water injection section does not indicate the average cooling rate to a temperature to be wound by the winder from injection start point.
[0134]
　Coiling temperature to 400 ° C. or higher 600 ° C. or less. This is a lower temperature than typical coiling temperature. The hot-rolled steel sheets produced in the conditions described above, by winding at this temperature range, the structure of the steel sheet, carbide fine ferrite can be dispersed bainite.
[0135]
　When the coiling temperature is below 400 ° C., transformed into untransformed a an austenitic stiff martensite before winding, during payout of the wound hot rolled steel sheet, cracks generated in the surface layer, the impact resistance since decreases, the winding temperature is set to 400 ° C. or higher. Preferably at 430 ℃ or more.
[0136]
　On the other hand, when the coiling temperature exceeds 600 ° C., a large pearlite lamellar spacing is produced, formed high thick acicular carbides of thermal stability, even after the annealing of the two-stage stepped, needle-like carbides There remains. During cold forging, cracking is the needle-shaped carbides starting, since the progress, the winding temperature is set to 600 ° C. or less. Preferably at 570 ℃ or less.
[0137]
　The hot-rolled steel sheet manufactured under the above conditions, after pickling, the annealing of 2-step type of holding at two temperature ranges applied. By performing annealing of the two-stage step-type hot-rolled steel sheet, to control the stability of the carbides, promote the formation of carbides to ferrite grain boundaries.
[0138]
　First, the technical idea of ​​the two-stage step-type annealing is described.
[0139]
　By performing the annealing of the first stage in a temperature range below Ac1 point, it causes the coarsening of carbides By concentrating additive metal element to increase the thermal stability of the carbide. Thereafter, the temperature was raised to a temperature range of not lower than Ac1 point austenite was produced in the tissue, the carbide in fine ferrite grains is dissolved in the austenite, leaving the coarse carbides in the austenite.
[0140]
　Subsequent slow cooling, austenite was transformed to ferrite, it will increase the carbon concentration in the austenite. By advancing the slow cooling, carbon atoms adsorbed on carbide remaining in the austenite, carbides and austenite, now cover the grain boundaries of the ferrite, finally, carbides in the ferrite grain boundary abundant tissue it is possible to form a. Therefore, tissue defining the present invention, it is apparent that not be formed only by a simple annealing.
[0141]
　Hereinafter, a description will be given of a specific annealing conditions.
[0142]
　[1-stage annealing]
　　heating rate to the annealing temperature: 30 ° C. / time or 0.99 ° C. / Time
　　annealing temperature: 650 ° C. or higher 720 ° C. or less
　　annealing temperature retention time: 60 hours or less than 3 hours
[0143]
　The heating rate up to the annealing temperature of the first stage and 30 ° C. / time or 0.99 ° C. / time or less. If the heating rate is less than 30 ° C. / time, since productivity takes time to raise the temperature is lowered, the heating rate is set to 30 ° C. / time or more. Preferably 40 ° C. / time or more.
[0144]
　On the other hand, if the heating rate exceeds 0.99 ° C. / time, temperature difference between the inner and the outer peripheral portion of the coil is increased, due to the thermal expansion difference occurs with scratches or burn, irregularities are generated on the surface of the steel sheet. During cold forging, cracking is the uneven starting, so deteriorating the impact resistance after reduction and carburizing and quenching and tempering the cold forgeability, the heating rate is less 0.99 ° C. / hour. Preferably at most 120 ° C. / hour.
[0145]
　1-stage annealing at annealing temperature (first stage annealing temperature) is set to 650 ° C. or higher 720 ° C. or less. The annealing temperature of the first stage is less than 650 ° C., the stability of the carbides becomes insufficient, the annealing in the second stage, since it is difficult to leave the carbide in austenite, annealing temperature of the first stage and 650 ℃ or more. Preferably at 670 ℃ or more.
[0146]
　On the other hand, if the annealing temperature exceeds 720 ° C., before the stability of the carbides increases, austenite is generated, since it is impossible to control the tissue changes described above, the annealing temperature is set to 720 ° C. or less. Preferably 700 ° C. or less.
[0147]
　(Retention time of the first stage) retention time in the annealing of the first stage is not more than 60 hours or more 3 hours. The holding time of the first stage is less than 3 hours, rather than stabilization of carbides is sufficiently in the annealing in the second stage, since it is difficult to leave the carbides, 3 hour hold time of the first stage and more. Preferably it is greater than or equal to 10 hours.
[0148]
　On the other hand, the retention time of the first stage is more than 60 hours, not be expected even more stability improvement of carbide, further, since the productivity is reduced, first stage holding time is set to less than 60 hours. Preferably not more than 50 hours.
[0149]
　[2-stage annealing]
　　heating rate to the annealing temperature: 1 ° C. / time or 80 ° C. / Time
　　annealing temperature: 725 ° C. or higher 790 ° C. or less
　　annealing temperature retention time: 50 hours or less than 3 hours
[0150]
　After completion of the holding in the annealing in the first stage, the hot-rolled steel sheet is heated below 80 ° C. / time or more heating rate 1 ° C. / time to annealing temperature. When cooled without annealing of the second stage, not ferrite grain size is large, it is impossible to obtain an ideal tissue.
[0151]
　In the annealing of the second stage, austenite is generated to grow from ferrite grain boundaries. By slowing the heating rate, it is possible to suppress the nucleation of austenite, in tissue obtained after slow cooling, it is possible to increase the grain boundary coverage carbide. Therefore, the heating rate in the annealing of the second stage is preferably small.
[0152]
　If the heating rate is less than 1 ° C. / time, since productivity takes time to raise the temperature is lowered, the heating rate is set to 1 ° C. / time or more. Preferably at 10 ° C. / time or more.
[0153]
　On the other hand, if the heating rate exceeds 80 ° C. / time, temperature difference between the inner and the outer peripheral portion of the coil is increased, due to the large thermal expansion difference by transformation, with scratches or burn occurs and irregularities on the surface of the steel sheet There is generated. During cold forging, the unevenness crack generated starting, so deteriorating the impact resistance after reduction and carburizing and quenching and tempering the cold forgeability, the heating rate is less 80 ° C. / hour.
[0154]
　2-stage annealing at the annealing temperature (the second stage annealing temperature) is set to 790 ° C. or less 725 ° C. or higher. The annealing temperature of the second stage is lower than 725 ° C., the amount of austenite becomes small, after the after the second stage annealing cooling, reduces the number ratio of carbides on the grain boundary of ferrite and ferrite grain size It becomes smaller. Therefore, the annealing temperature in the second stage to 725 ° C. or higher. Preferably at 735 ℃ or more.
[0155]
　On the other hand, if the annealing temperature of the second stage is higher than 790 ° C., it becomes difficult to leave carbides in the austenite, since it is difficult to control the tissue changes described above, the annealing temperature of the second stage 790 ° C. or less to. Preferably at 780 ℃ or less.
[0156]
　(Retention time in the second stage) retention time in the annealing in the second stage is not more than 50 hours or more 1 hour. If the holding time of the second stage is less than 1 hour, less the amount of austenite and, dissolution of carbides in the ferrite grains is not sufficient, to increase the number ratio of carbides on the grain boundary of ferrite It becomes difficult, and since the ferrite grain size decreases, the second stage of the retention time is 1 hour or longer. Preferably it is greater than or equal to 5 hours.
[0157]
　On the other hand, the retention time in the second stage is more than 50 hours, so it is difficult to leave the carbide in austenite, second-stage retention time shall be 50 hours or less. Preferably not more than 45 hours.
[0158]
　Cooling after annealing]
　　Cooling stop temperature: 650 ° C.
　　Cooling rate: 1 ° C. / time or longer 100 ° C. / time or less
[0159]
　After holding at the annealing in the second stage is finished, the hot-rolled steel sheet after annealing, to 650 ° C., then slowly cooled 100 ° C. / Time 1 ° C. / time or more of the following cooling rates. When the stop temperature of the gradual cooling is more than 650 ° C., the austenite un transformation by the cooling rate in excess of 100 ° C. / time to subsequent room temperature, transformed into pearlite or bainite, increased hardness, cold forgeability since reduced, the cooling stop temperature is set to 650 ° C..
[0160]
　The austenite generated in the annealing in the second stage is cooled, with to transform to ferrite, in order to adsorb the carbon to carbide remaining in the austenite, the cooling rate is slower is preferable. When the cooling rate is less than 1 ° C. / time, increases the time required for cooling, since productivity is lowered, the cooling rate is set to 1 ° C. / time or more. Preferably at 10 ° C. / time or more.
[0161]
　On the other hand, if the cooling rate exceeds 100 ° C. / time, austenite transforms to pearlite, increases the hardness of the steel sheet, so deteriorating the impact resistance after reduction and carburizing and quenching and tempering the cold forgeability, the cooling rate is not more than 100 ° C. / hour. Preferably from 90 ° C. / hour.
[0162]
　Here, the cooling stop temperature is that the temperature to be controlled by the cooling rate, by performing the cooling to 650 ° C. Speed ​​1 ° C. / time or longer 100 ° C. / time less cooling, to 650 ° C. or less the cooling is not particularly limited.
[0163]
　Incidentally, the atmosphere of the annealing is not limited to a specific atmosphere. For example, an atmosphere of 95% nitrogen, 95% or more an atmosphere of hydrogen, and may be any of air atmosphere.
[0164]
　As described above, the hot rolling conditions and annealing conditions consistently manage present invention, according to the production method for performing tissue control of the steel plate, the diaphragm, between the superior cold to the cold forging the thickened molding was Kumia' exhibits forgeability can further produce a low carbon steel sheet excellent in impact resistance after carburizing quenching and tempering.
Example
[0165]
　Next is a description of examples, the level of embodiments is an example of employing the execution conditions for confirming the workability and effects of the present invention, the present invention is limited to this single example of conditions not. The present invention does not depart from the gist of the present invention, as long as to reach the present invention purpose is as it can employ various conditions.
[0166]
　Continuous casting slab having the component composition shown in Table 1 (steel ingot) was heated for 1.8 hours at 1240 ° C., and subjected to hot rolling. Exit finish hot rolled at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C. / sec on ROT, wound at 510 ° C., to produce a hot-rolled coil having a thickness of 5.2 mm.
[0167]
[Table 1]

[0168]
　The hot rolled coil was pickled, it was charged with the coil in a box type annealing furnace, after controlling the atmosphere of 95% hydrogen -5% nitrogen, and heated at a heating rate of 100 ° C. / time to 705 ° C. from room was uniform temperature distribution in the coil and held for 36 hours at 705 ° C.. Then, the mixture was heated to 760 ° C. at a heating rate of 5 ° C. / time, further, after holding for 10 hours at 760 ° C., until 650 ° C. and cooled at a cooling rate of 10 ° C. / time, and then cooled in the furnace to room temperature , to prepare a sample for property evaluation.
[0169]
　Samples of tissues were observed by the method described above, crack length present in the sample after the cold forging was measured by the method described above.
[0170]
　Thickened molded carburizing of the sample was carried out by gas carburizing. From the furnace atmospheric gas to diffuse carbon into the steel interior through the sample surface, in a furnace with controlled carbon potential of 0.5 wt% C, performs a process of holding 120 minutes at 940 ° C., then at room temperature the sample was cooled in the furnace to.
[0171]
　Subsequently, after heating to 840 ° C. from room temperature and the retention of 20 minutes, quenching it was in 60 ° C. oil. Quenching the sample was subjected to tempering process of cooling after 60 min holding at 170 ° C., to produce a carburizing quenching and tempering samples.
[0172]
　The impact resistance of the carburized quenching and tempering samples were evaluated by drop weight test. Figure 2 shows an overview of drop weight test for evaluating the impact resistance of the samples subjected to carburizing quenching and tempering schematically. The carburizing quenching and tempering the cup-shaped cup bottom of the sample 4 having undergone fixed with jig, the cup sides, weight 2kg falling weight (upper width: 50 mm, lower side width: 10 mm, height: 80 mm, length: 110 mm) were allowed to free fall from 4m away top, the vertical wall portion of the sample 4, shocked about 80 J, to observe the presence or absence of cracks in the sample, to evaluate the impact resistance.
[0173]
　A result of free fall, for the sample did not show a crack or destruction, with a rating of excellent "OK" to the impact resistance, for a sample that was seen cracking and destruction, poor impact resistance "NG" of It was scored.
[0174]
　Table 2, the maximum of the samples produced, carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, the ratio of the number of carbides of the ferrite grain boundaries to the number of carbide in ferrite grains, for the thickness of the vertical wall portion ratio of crack length, and shows the evaluation results and measurement results of the impact resistance.
[0175]
[Table 2]

[0176]
　As shown in Table 2, inventive steels A-1, B-1, C-1, D-1, E-1, F-1, G-1, H-1, I-1, J-1 and, , K-1 are both greater than the ratio of the number of carbides 1 ferrite grain boundaries to the number of carbide in ferrite grains, Vickers hardness is less 180HV than 100 HV, carburizing quenching and tempering the cold forgeability It has excellent impact resistance after.
[0177]
　In contrast, the comparative steel L-1 has a low C content, because the hardness before cold forging is less than 100 HV, low cold forgeability. Comparative Steel M-1, P-1, and, Z-1 is, P, Al, excessively containing N, at the second stage annealing, since segregation to gamma / alpha interface is large, the grain boundary formation of carbides is suppressed.
[0178]
　Comparative Steel S-1 contains excess of Si, for ferrite ductility is low, a low cold forgeability. Comparative Steel N-1 and T-1, respectively, Mo, because it contains an excess of Cr, carbides finely dispersed in the ferrite grains and the hardness exceeds 180 Hv. Comparative Steel Q-1, because the excessively containing Mn, is significantly lower impact properties after carburizing quenching and tempering.
[0179]
　Comparative Steel O-1 is less Cr amount, because the surface layer of the austenite grains is abnormally coarsened during carburizing, a low impact resistance. Comparative Steel R-1, because the excessively containing S, coarse MnS is generated in the steel, low cold forgeability. Comparative Steel U-1, since excessively containing C, coarse carbides are generated within the thickness increase of the steel, for coarse carbides after carburizing quenching remained, low impact resistance.
[0180]
　Comparative Steel V-1 have less Mn amount, since it is difficult to increase the stability of carbides, low impact resistance after cold forgeability and carburizing quenching and tempering. Comparative Steel W-1 and X-1 is, O, for excessively containing Ti, oxides present in the ferrite grain, TiC becomes the carbide formation site in slow cooling after 2 phase region annealing, the grain boundary generation of carbides is suppressed in, low cold forgeability. Comparative Steel Y-1, since excessively containing B, and a low cold forgeability.
[0181]
　Subsequently, in order to examine the influence of the manufacturing conditions, A shown in Table 1, B, C, D, E, F, G, H, I, J, and, a slab having the component composition of K, shown in Table 3 at hot rolling conditions and annealing conditions, to prepare a hot-rolled sheet annealing sample thickness 5.2 mm.
[0182]
　Table 4, for the sample prepared, carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, the ferrite grain boundaries to the number of carbide in ferrite grains ratio of the number of carbides, for the thickness of the vertical wall portion percentage of maximum crack length, and shows the evaluation results and measurement results of the impact resistance.
[0183]
[table 3]

[0184]
[Table 4]

[0185]
　Comparative Steel E-3 has low finishing hot-rolling temperature, low productivity rolling load is increased. Comparative Steel D-2 has high finishing hot-rolling temperature, since scale defects on the surface of the steel sheet is produced, when subjected to abrasion test after quenching and tempering, cracking and peeling occur a scale defects starting, wear characteristics were lowered. Comparative Steel F-2 is, ROT (Run Out Table) the cooling rate is slow at, has led to occurrence of reduction and scale defects productivity.
[0186]
　Comparative steel C-4 is a cooling rate 100 ° C. / sec in ROT, by the outermost surface layer portion of the steel sheet is excessively cooled, fine cracks on the outermost surface layer portion was formed. Comparative Steel C-2 is low coiling temperature, low temperature transformation structure such as bainite or martensite is often produced by embrittlement, cracking during hot rolling coil payout is frequently, productivity is lowered. Further, wear resistance in the sample taken from cracking strip is low.
[0187]
　Comparative Steel G-2 has a high coiling temperature, with thick pearlite generates a lamellar spacing in the hot-rolled tissue, needle-like thermal stability of the coarse carbide is high, after the annealing of 2-step type also, since the carbide remains in the steel sheet, low machinability. Comparative Steel H-4, since the heating rate in the first-stage annealing annealing of the two-stage stepped slow, the productivity is low.
[0188]
　Comparative Steel E-3, because the heating rate in the annealing of the first stage is high, the temperature difference between the inside and the inner peripheral portion of the coil is increased, scratches and seizure caused by thermal expansion difference occurs, hardening when subjected to the evaluation test of the antiwear properties after tempering, cracking and peeling occurred from the flaw portion, wear resistance is lowered.
[0189]
　It Comparative Steel G-4 is the first stage of holding at the annealing temperature (annealing temperature) is low, coarse processing carbide below Ac1 point is insufficient, the thermal stability of the carbides is insufficient Accordingly, carbides remaining in the annealing in the second stage is reduced, it can not be suppressed pearlite transformation in tissue after annealing, low machinability.
[0190]
　Comparative Steel D-4 has a high holding temperature of the annealing in the first stage (annealing temperature), austenite formed during annealing, it is not possible to increase the stability of carbides, pearlite formed after annealing, Vickers hardness beyond Saga 180 Hv, less machinability. Comparative Steel J-4 has a short retention time in the annealing in the first stage, it is impossible to increase the stability of carbides, low machinability.
[0191]
　Comparative Steel F-2 has a long retention time in the annealing in the first stage, in addition to low productivity, seizure flaws occurred, a low wear resistance. Comparative Steel B-4, since the heating rate in the second-stage annealing annealing of the two-stage stepped slow, the productivity is low. Comparative steel A-3, since the heating rate in the annealing of the second stage is high, the temperature difference between the inner and the outer peripheral portion of the coil is increased, it is generated scratches and seizure due to large thermal expansion difference by transformation, low wear resistance after quenching and tempering.
[0192]
　Comparative Steel K-2, the lower holding temperature in the annealing of the second stage (annealing temperature), less the amount of austenite, since it is not possible to increase the ratio of the number of carbides in the ferrite grain boundary, a low machinability. Comparative steel C-4, the holding temperature in the annealing of the second stage (annealing temperature) is high, since the dissolution of carbides is promoted during annealing, it is difficult to form a grain boundary carbides after annealing, further perlite generated, Vickers hardness exceeds the 180 Hv, less machinability.
[0193]
　Comparative Steel J-3 has a long retention time in the annealing in the second stage, since the dissolution of carbides is promoted, low machinability. Comparative Steel D-3, the slow cooling rate to 650 ° C. from the second-stage annealing, with low productivity, tissues generate coarse carbides after slow cooling, coarse carbides at the time of cold forging cracking occurs as a starting point, cold forging of which was lowered. Comparative Steel I-3, the fast cooling rate from the second-stage annealing to 650 ° C., since the hardness happening pearlite transformation is increased during the cooling, is low cold forgeability.
[0194]
　Next, in order to examine the allowable content of the other element, after Tables 5 and 6 for 1.8 hours continuous casting slab having the composition shown in (continuation of Table 5) and (steel ingot) at 1240 ° C. heating, heat They were subjected during rolling. Exit finish hot rolled at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C. / sec on ROT, wound at 510 ° C., to produce a hot-rolled coil having a thickness of 5.2 mm.
[0195]
[table 5]

[0196]
[Table 6]

[0197]
　The hot rolled coil was pickled, it was charged with hot rolled coil in a box type annealing furnace, after controlling the atmosphere of 95% hydrogen -5% nitrogen, at a heating rate of 100 ° C. / time to 705 ° C. from room heating, and held for 36 hours at 705 ° C. to equalize the temperature distribution in the coil, then heated to 760 ° C. at a heating rate of 5 ° C. / time, further, after holding for 10 hours at 760 ° C., until 650 ° C. It was cooled at a cooling rate of 10 ° C. / time, then furnace cooling to room temperature, to prepare a sample for property evaluation.
[0198]
　Incidentally, a sample of tissue, was observed by the method described above, crack length present in the sample after the cold forging was measured by the method described above.
[0199]
　Table 7, the maximum in the samples prepared, carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, the ratio of the number of carbides of the ferrite grain boundaries to the number of carbide in ferrite grains, for the thickness of the vertical wall portion ratio of crack length, and shows the evaluation results and measurement results of the impact resistance.
[0200]
[Table 7]

[0201]
　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 both the ratio of the number of ferrite grain boundary carbide to the number of carbide in ferrite grains There than 1, Vickers hardness of not more than 180HV than 100 HV, is excellent in impact resistance after carburizing quenching and tempering the cold forgeability.
[0202]
　In contrast, the comparative steels AR-1, AS-1, AW-1, AZ-1, BB-1, and, BF-1 is contained excessively, respectively, La, As, Cu, Ni, Sb, Ce, and, since the segregation amount of the gamma / alpha interface during the second stage annealing increases, generation of carbides at the grain boundaries is suppressed. Comparative Steel BG-1 contains excess of Si, for ferrite ductility is low, a low cold forgeability.
[0203]
　Comparative Steel AT-1, AV-1, BA-1, BC-1, BH-1, and, BJ-1, respectively, Mo, Nb, Cr, Ta, W, for excessively containing V, ferrite carbides finely dispersed in grains, and the hardness is greater than 180 Hv. Comparative Steel BF-1, since excessively containing Mn, is significantly lower impact properties after carburizing quenching and tempering.
[0204]
　Comparative Steel AU-1, AX-1, AY-1, and, BE-1, respectively, Zr, Ca, Mg, containing an excess of Y, coarse oxides or non-metallic inclusions in the steel product to, cracking is starting from the coarse oxide or coarse nonmetallic inclusions in the cold forging, cold forging property deteriorate. Comparative Steel BD-1 is excessively contained Sn, ferrite is brittle, low cold forgeability. Comparative Steel BK-1, since excessively containing C, to produce coarse carbides within the thickness increase of the steel, coarse carbides after carburizing quenching remains, impact resistance is lowered.
[0205]
　Subsequently, in order to examine the influence of the manufacturing conditions, AA shown in Table 5, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP, and, AQ slabs having the chemical composition, hot-rolled condition and the annealing condition shown in Table 8, to prepare a hot-rolled sheet annealing sample thickness 5.2 mm.
[0206]
[Table 8]

[0207]
　Table 9, the maximum in the sample prepared, carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, the ratio of the number of ferrite grain boundary carbide to the number of carbide in ferrite grains, for the thickness of the vertical wall portion ratio of crack length, and shows the evaluation results and measurement results of the impact resistance.
[0208]
[Table 9]

[0209]
　Comparative Steel AC-2, the finishing hot-rolling temperature is low, the productivity is low. Comparative Steel AN-4 is a high finishing hot-rolling temperature, and generation of scale defects on the steel sheet surface, when an impact load is applied after cold forging and carburizing quenching and tempering, cracks generated from the flaw portion, impact characteristics were lowered.
[0210]
　Invention steel AB-3, since the cooling rate in the ROT is slow, led to derivation of reduction and scale defects productivity. Invention steel AJ-3 and AD-4 is a 100 ° C. / sec cooling rate in ROT, the outermost surface layer portion is excessively cooled steel sheet, fine cracks on the outermost surface layer portion was formed.
[0211]
　Comparative Steel AN-3 is lower coiling temperature, low temperature transformation structure such as bainite or martensite is often produced by embrittlement, the productivity was reduced cracks frequently occur during the hot-rolled coil payout. Moreover, impact resistance after cold forging and carburizing quenching and tempering in the samples taken from the cracked pieces was inferior.
[0212]
　Comparative Steel AH-3 has a high coiling temperature, with thick pearlite generates a lamellar spacing in the hot-rolled tissue, needle-like thermal stability of the coarse carbide is high, even after the 2-step type annealing , since the carbide remains in the steel sheet, low cold forgeability.
[0213]
　Comparative Steel AF-4, since the heating rate in the first-stage annealing annealing of the two-stage stepped slow, the productivity is low. Comparative Steel AG-2, since the heating rate in the annealing of the first stage is high, the temperature difference between the inner and the outer peripheral portion of the coil is increased, scratches and seizure caused by thermal expansion difference occurs, cold impact resistance after forging and carburizing quenching and tempering is lowered.
[0214]
　Comparative Steel AA-2 is the first stage of holding at the annealing temperature (annealing temperature) is low, insufficient coarsening process carbide below Ac1 point, the thermal stability of the carbides becomes insufficient, 2-stage eye carbides remaining decreases during annealing can not suppress pearlite transformation in structure after annealing, cold forging property deteriorate.
[0215]
　Comparative Steel AM-3 has a high first stage holding temperature (annealing temperature), austenite formed during the anneal, it is impossible to increase the stability of carbides, cold forgeability and resistance after carburizing quenching and tempering impact properties was reduced. Comparative Steel AF-2 has a short retention time in the annealing in the first stage, it is impossible to increase the stability of the carbides is less cold forgeability. Comparative Steel AO-4 has a long retention time in the annealing in the first stage, the productivity is low.
[0216]
　Comparative Steel AP-4, since the heating rate in the second-stage annealing annealing of the two-stage stepped slow, the productivity is low. Comparative Steel AI-3, since the heating rate in the annealing of the second stage is high, the temperature difference between the coil inner and the outer peripheral portion becomes large, scratches and seizure due to large thermal expansion difference due to transformation occurs, carburizing when an impact load is applied after quenching and tempering, cracking is from 該疵 part, impact resistance is lowered.
[0217]
　Comparative Steel AL-3 is lower holding temperature in the annealing of the second stage (annealing temperature), less the amount of austenite, can not increase the number ratio of carbides in the ferrite grain boundary, cold forgeability reduced. Comparative Steel AD-2 has high holding temperature in the annealing of the second stage (annealing temperature), because the dissolution of carbides is promoted during annealing, it is difficult to produce a grain boundary carbides after annealing, cold forgeability and impact properties after carburizing quenching and tempering is lowered.
[0218]
　Comparative Steel AJ-4 has a long retention time in the annealing in the second stage, since the dissolution of carbides is promoted, low cold forgeability. Comparative Steel AQ-3, the slow cooling rate to 650 ° C. from the second-stage annealing, with low productivity, tissue coarse carbides are generated after slow cooling, during cold forging, coarse carbides cracking is starting a cold forging property deteriorate. Comparative Steel AP-2, the fast cooling rate from the second-stage annealing to 650 ° C., occurs pearlite transformation during cooling, because the hardness is increased, cold forging property deteriorate.
[0219]
　Here, FIG. 3 shows the ratio of the number of grain boundary carbides to the number of grain carbide, the relationship between the crack length and the impact resistance after carburizing quenching and tempering of the cold forging test piece.
[0220]
　3, when the number ratio (= the number of number / intragranular carbide grain boundary carbides) exceeds 1, it is possible to suppress the ratio of the crack length that is introduced by cold forging, excellent after carburizing quenching and tempering impact resistance seen that obtained with.
[0221]
　Also, it is shown in FIG. 4, the ratio of the number of grain boundary carbides to the number of grain carbide, another relationship crack length and impact resistance after carburizing quenching and tempering of the cold forging test piece. 4, even in the steel sheet obtained by adding an additive element, illustrates that it is possible to suppress the crack length.
[0222]
　From Figure 4, the steel sheet in the case of adding an element proper range also, the number ratio (= the number of number / intragranular carbide grain boundary carbides) When more than 1, the crack length that is introduced by cold forging it is possible to suppress the ratio, it can be seen that excellent impact resistance after carburizing quenching and tempering can be obtained.
Industrial Applicability
[0223]
　As described above, according to the present invention, it is possible to provide a low carbon steel sheet and a manufacturing method thereof excellent in impact resistance after cold forgeability and carburizing quenching and tempering. Steel sheet of the present invention, for example, since it is suitable as a material when molded by cold forging a plate molding or the like to obtain a component such as high cycle gear, the present invention is, as it has high industrial applicability is there.
DESCRIPTION OF SYMBOLS
[0224]
　1 disc-shaped test material
　2 cup-shaped test material
　-out 3 Cracks
　4 Sample
　5 falling weight
　maximum length of L cra

claims

[Claim 1]Component composition, bymass%,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: than
　　0.020%, 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%
　　hereinafter, Ca: 0.050% or
　　less, Y: 0.050% or
　　less, Zr: 0.050% or
　　less, La: 0.050% or less , And
　　Ce: 0.050%
comprises, a low carbon steel sheet which is a balance of Fe and impurities,
　metal structure of the low-carbon steel sheet,
　carbide particle size of 0.4 ~ 2.0 .mu.m,
　perlite area rate 6 % or less, and
　ferrite grains in the ratio of the number of ferrite grain boundary carbide to the number of carbide is 1 greater than
a filled,
　Vickers hardness of the low carbon steel is less than 180HV than 100HV
steel sheet, characterized in that.
[Claim 2]
　A method of manufacturing a steel sheet according to claim 1,
　is subjected to between completing hot rolling finish hot rolled in a temperature range of 650 ° C. or higher 950 ° C. or less slab component composition according to claim 1 and hot-rolled steel sheet,
　the hot-rolled steel sheet was wound at 400 ° C. or higher 600 ° C. or less,
　wound subjected to pickling hot-rolled steel sheet, pickled steel sheet below 30 ° C. / time or 0.99 ° C. / time hot rolling at a heating rate, and heated to a annealing temperature 720 ° C. 650 ° C. or higher, annealed in the first stage to 3 hours or more 60 hours or less, then
　hot-rolled steel sheet 1 ° C. / time or 80 ° C. / time or less in the heating rate and heated to an annealing temperature below 790 ° C. 725 ° C. or higher, annealed in the second stage to hold 50 hours or less than 3 hours, the hot-rolled steel sheet after annealing, 1 ° C. / time or longer 100 ° C. It cooled to 650 ° C. / time less cooling rate
, characterized in that Method of manufacturing the plate.