US20180135146A1

US20180135146A1 - Steel sheet and method of production of same

Info

Publication number: US20180135146A1
Application number: US15/576,682
Authority: US
Inventors: Ken Takata; Kazuo Hikida; Kengo Takeda; Motonori Hashimoto
Original assignee: Nippon Steel and Sumitomo Metal Corp
Current assignee: Nippon Steel Corp
Priority date: 2015-05-26
Filing date: 2016-05-26
Publication date: 2018-05-17
Also published as: KR20170138510A; CN107614726A; CN107614726B; JP6119924B1; EP3305930A1; JPWO2016190396A1; EP3305930A4; MX2017014938A; BR112017024957A2; TWI612154B; WO2016190396A1; TW201704501A; KR102029566B1

Abstract

A steel sheet improved in hardenability and material formability having a predetermined chemical composition, characterized in that, in the metal structure of the steel sheet, an average grain size of carbides is 0.4 μm to 2.0 μm, an area ratio of pearlite is 6% or less, when a number of carbides in ferrite grains is A and a number of carbides at ferrite grain boundaries is B, B/A>l, and when an X-ray diffraction intensity at {211}<011>at a plane of a part of ½ sheet thickness of the steel sheet is denoted by “I1” and an X-ray diffraction intensity at {100}<011>is denoted by “I0”, I1/I0<1 is satisfied, and the steel sheet has a Vickers hardness of 100 HV to 150 HV.

Description

TECHNICAL FIELD

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

BACKGROUND ART

Steel sheet containing, by mass %, carbon in an amount of 0.1 to 0.7% is being used as a material for production of gears, clutches, and other drive system parts of automobiles by being used press-formed, enlarged in holes, bent, drawn, thickened, and thinned and cold forged by combinations of the same from a blank. The strength of such parts is secured by quenching and tempering, so a high hardenability is demanded from steel sheet.
Furthermore, a high formability in the cold state is demanded from steel sheet used as a material for such drive system parts. Parts are mainly formed by drawing and/or thickening. In forming parts, the biggest factor governing the material formability is the plastic anisotropy. Improvement of the plastic anisotropy in steel sheet is necessary for application of steel sheet to the formation of parts.
Several proposals have been made up to now for the hardenability demanded and formability improved in plastic anisotropy. The following patent literature discloses steel sheet excellent in cold forgeability and impact resistance characteristic.
For example, PLT 1 discloses, as steel for machine structural use improving toughness by suppressing coarsening of crystal grains in carburization heat treatment, steel for machine structural use containing, by mass %, C: 0.10 to 0.30%, Si: 0.05 to 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 to 0.050%, Nb: 0.02 to 0.10%, and N: 0.0300% or less and having a balance of Fe and unavoidable impurities, having a structure before cold working comprised of ferrite and pearlite structures, and having an average value of ferrite grain size of 15 μm or more.
PLT 2 discloses, as steel excellent in cold workability and carburizing and quenching ability, steel containing C: 0.15 to 0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, and B: 0.005 to 0.050%, having a balance of Fe and unavoidable impurities, and having a structure mainly comprised of ferrite phases and graphite phases.
PLT 3 discloses a steel material for carburized bevel gear use excellent in impact strength, a high toughness carburized bevel gear, and a method of production of the same.
PLT 4 discloses, for a part produced by spheroidal annealing, then a cold forging and a carburizing, quenching, and tempering process, steel for carburized part use having excellent workability while suppressing coarsening of crystal grains even with subsequent carburization and having an excellent impact resistance characteristic and impact fatigue resistance characteristic.
PLT 5 discloses as cold tool steel for plasma carburization use a steel containing C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn: 0.05 to 1.50%, and V: 1.8 to 6.0%, further containing one or more of Ni: 0.10 to 2.50%, Cr: 0.1 to 2.0%, and Mo: 3.0% or less, and having a balance of Fe and unavoidable impurities.
On the other hand, there have been the following proposals for improvement of the formability, that is, the improvement of plastic anisotropy.
For example, PLT 6 proposes prescribing the carbide grain size and spheroidization rate in steel containing C: 0.25 to 0.75% and improving the in-plane anisotropy by the cold rolling rate and box annealing conditions, the coiling temperature in hot rolling, and provisions on the texture so as to limit the “r” value and Δr.
PLTs 7 and 8 propose to prescribe the heating and annealing conditions of a hot rolled material between stands of a finish rolling machine so as to reduce the Δr value and improve the in-plane anisotropy. PLT 8 proposes steel sheet reduced in in-plane anisotropy by prescribing hot rolling during which performing finish rolling at a temperature of the Ar3 point or more and coiling at 500 to 630° C.

CITATION LIST

Patent Literature

PLT 1: Japanese Patent Publication No. 2013-040376A
PLT 2: Japanese Patent Publication No. 06-116679A
PLT 3: Japanese Patent Publication No. 09-201644A
PLT 4: Japanese Patent Publication No. 2006-213951A
PLT 5: Japanese Patent Publication No. 10-158780A
PLT 6: Japanese Patent Publication No. 2000-328172A
PLT 7: Japanese Patent Publication No. 2001-073076A
PLT 8: Japanese Patent Publication No. 2001-073077A

SUMMARY OF INVENTION

Technical Problem

The above patent literature proposed improvement of the in-plane anisotropy, but did not propose the provision of the strength demanded from the part, that is, the hardenability.
The present invention was made in consideration of the above situation in the prior art and has as its object the provision of steel sheet improved in hardenability and material formability, in particular, optimal for obtaining a gear or other part by thickening or other cold forging, and a method of production of the same.

Solution to Problem

To solve the above problem and obtain steel sheet suitable for the material of a drive system part etc., it can be understood that in steel sheet containing the C necessary for raising the hardenability, it is sufficient to increase the grain size of the ferrite, spheroidize the carbides (mainly cementite) by a suitable grain size, and decrease the pearlite structures. This is due to the following reasons.
Ferrite phases are low in hardness and high in ductility. Therefore, in a structure mainly comprised of ferrite, it becomes possible to increase the grain size so as to raise the material formability.
Carbides, by being made to suitably disperse in the metal structure, can maintain the material formability while imparting an excellent wear resistance and rolling fatigue characteristic, so are structures essential for drive system parts. Further, the carbides in the steel sheet are strong particles obstructing slip. By forming carbides at the ferrite grain boundaries, it is possible to prevent propagation of slip exceeding the crystal grain boundaries and suppress the formation of shear zones. The cold forgeability is improved and, simultaneously, the formability of steel sheet is also improved.
However, cementite is a hard, brittle structure. If a laminar structure with ferrite present, that is, in the state of pearlite, the steel becomes hard and brittle, so it has to be present in a spheroidal form. If considering the cold forgeability and the occurrence of fractures at the time of forging, its grain size has to be a suitable range.
However, no method of production for realizing the above structure has been disclosed up to now. Therefore, the inventors intensively researched a method of production for realizing the above structure.
As a result, they discovered the following: To make the metal structure of the steel sheet after coiling after hot rolling a bainite structure of fine pearlite or fine ferrite with small lamellar spacing in which cementite is dispersed, the steel sheet is coiled at a relatively low temperature (400° C. to 550° C.). By coiling at a relatively low temperature, the cementite dispersed in the ferrite also easily becomes spheroidal. Next, the cementite is partially made spheroidal by annealing at a temperature just under the Ac1 point as first stage annealing. Next, as second stage annealing, part of the ferrite grains is left while part is transformed to austenite by annealing at a temperature between the Ac1 point and Ac3 point (so-called dual phase region of ferrite and austenite). By then making the remaining ferrite grains grow while slowly cooling the steel while using these as nuclei to transform the austenite to ferrite, it is possible to obtain large ferrite phases and make cementite precipitate at the grain boundaries to realize the above structure.
That is, the inventors found that it is difficult to realize a method of production of steel sheet satisfying both hardenability and formability even if adjusting the heat rolling conditions, annealing conditions, etc. separately and that it is possible to realize this by optimization by a so-called integrated process of hot rolling, annealing, etc.
Further, they found that for improvement of the drawability at the time of cold forming, reduction of the plastic anisotropy is necessary and, to improve this, adjustment of the hot rolling conditions is important.
The present invention was made based on these findings and has as its gist the following:
(1) A steel sheet consisting of, by mass %, C: 0.10 to 0.70%, Si: 0.01 to 0.30%, Mn: 0.30 to 3.00%, Al: 0.001 to 0.10%, Cr: 0.010 to 0.50%, Mo: 0.0010 to 0.50%, B: 0.0004 to 0.01%, Ti: 0.001 to 0.10%, P: 0.02% or less,
S: 0.01% or less, N: 0.0200% or less, 0: 0.0200% or less, Sn: 0.05% or less, Sb: 0.05% or less, As: 0.05% 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: 0.05% or less, Ca: 0.05% or less, Y: 0.05% or less, Zr: 0.05% or less, La: 0.05% or less, and Ce: 0.05% or less and a balance of Fe and unavoidable impurities, wherein the metal structure of the steel sheet includes carbide having an average grain size of 0.4 μm to 2.0 μm, perlite having an area ratio of 6% or less, and ferrite wherein a ratio of a number of carbides at ferrite grain boundaries to a number of carbides in ferrite grains of over 1; and I1/I0<1 being satisfied when an X-ray diffraction intensity at {211}<011>at a plane of a part of ½ sheet thickness of the steel sheet is denoted by “I1” and an X-ray diffraction intensity at {100}<011>is denoted by “I0”, the steel sheet having a Vickers hardness of 100 HV to 150 HV.
(2) A method of production for producing steel sheet according to (1) comprising hot rolling a steel slab of a chemical composition according to (1) with finish rolling temperature between 820° C. and 950° C., to obtain hot rolled steel sheet; coiling the hot rolled steel sheet at 400° C. to 550° C.; pickling the coiled hot rolled steel sheet; heating the pickled hot rolled steel sheet to an annealing temperature of 650° C. to 720° C. by a heating rate of 30° C/hour to 150° C/hour and holding the steel sheet for 3 hours to 60 hours as a first stage of annealing; next, heating the hot rolled steel sheet to an annealing temperature of 725° C. to 790° C. by a heating rate of 1° C/hour to 80° C/hour and holding the steel sheet for 3 hours to less than 10 hours as a second stage of annealing; and, next, cooling the annealed hot rolled steel sheet to 650° C. by a cooling rate of 1° C/hour to 100° C/hour.

Advantageous Effects of Invention

According to the present invention, it is possible to provide steel sheet excellent in hardenability and material formability, in particular, optimal for obtaining a gear or other part by forming by thickening or other cold forging, and a method of production of the same.

DESCRIPTION OF EMBODIMENTS

Below, the present invention will be explained in detail. First, the reasons for limitation of the chemical composition of the steel sheet of the present invention will be explained. Here, the “%” according to the chemical composition means “mass %”.
C: 0.10 to 0.70%
C is an element forming carbides and effective for strengthening the steel and refining the ferrite grains. To suppress the formation of a matte surface in cold working and secure surface beauty of a cold forged part, suppression of coarsening of the ferrite grain size is necessary.
If C is less than 0.10%, the carbides become insufficient in volume fraction and coarsening of the carbides during annealing can no longer be suppressed, so C is made 0.10% or more. Preferably it is 0.14% or more. On the other hand, if the content of C increases, the carbides increase in volume fraction, cracks are formed acting as starting points of breakage at the time of an instantaneous load, and there is the fear that the formability and impact resistance characteristic will fall. If making this drop as small as possible, C is made 0.40% or less. Preferably it is 0.38% or less.
On the other hand, if the carbides increase in volume fraction and the strength rises, the fatigue characteristic is improved, so when improving the fatigue characteristic, C is made over 0.40%. Preferably it is 0.44% or more. If C is over 0.70%, a large amount of cracks forming starting points of breakage are formed and the fatigue characteristic conversely falls, so C is made 0.70% or less. Preferably it is 0.66% or less.
Si: 0.01 to 0.30%
Si is an element which acts as a deoxidizing agent and further has an effect on the form of the carbides and contributes to the improvement of the material formability. To obtain the deoxidizing effect, Si is made 0.01% or more. Preferably it is 0.07% or more.
If Si is over 0.30%, due to solution strengthening of the ferrite, the hardness rises and the ductility falls, fractures easily occur at the time of cold forging, and the formability at the time of cold forging and the impact resistance characteristic after carburization, quenching, and temperature falls, so Si is made 0.30% or less. Preferably it is 0.28% or less.
Mn: 0.30 to 3.00%
Mn is an element controlling the form of carbides in two-stage annealing. If less than 0.30%, in the gradual cooling after second stage annealing, it becomes difficult to form carbides at the ferrite grain boundaries, so Mn is made 0.30% or more. Preferably it is 0.40% or more.
If Mn is over 1.00%, after carburization, quenching, and tempering, the toughness falls, but on the other hand, the strength rises. When trying to keep down the drop in toughness after carburization, quenching, and tempering as much as possible Mn is made 1.00% or less. Preferably it is 0.96% or less.
When trying to raise the strength, Mn is made over 1.00%. Preferably it is 1.10% or more. If Mn is over 3.00%, after carburization, quenching, and tempering, the toughness remarkably falls, so Mn is made 3.00% or less. Preferably it is 2.70% or less.
Al: 0.001 to 0.10%
Al is an element which acts as a deoxidizing agent and stabilizes ferrite. If less than 0.001%, the effect of addition is not sufficiently obtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.
On the other hand, if Al is over 0.10%, the number of carbides at the ferrite grain boundaries decreases and the formability falls, so Al is made 0.10% or less. Preferably it is 0.09% or less.
Cr: 0.010 to 0.50%
Cr is an element effective for stabilization of carbides at the time of heat treatment. If less than 0.010%, it becomes difficult to cause carbides to remain at the time of carburization, coarsening of the austenite grain size at the surface layer is invited, and the strength drops, so Cr is made 0.010% or more. Preferably it is 0.050% or more.
On the other hand, if Cr is over 0.50%, the amount of Cr concentrating at the carbides increases and a large amount of fine carbides remain in the austenite phases produced by the two-stage annealing, carbides remain in the ferrite grains after gradual cooling inviting an increase in the hardness and a decrease in the number of carbides at the ferrite grain boundaries fall, and the formability falls, so Cr is made 0.50% or less. Preferably it is 0.40% or less.
Mo: 0.001 to 0.50%
Mo, like Mn and Cr, is an element effective for control of the form of carbides. If less than 0.001%, the effect of addition is not obtained, so Mo is made 0.001% or more. Preferably it is 0.005% or more.
On the other hand, if over 0.50%, Mo concentrates at the carbides, stable carbides increase even in the austenite phases, carbides remain inside the ferrite grains after gradual cooling inviting an increase in the hardness and a decrease in the number of carbides at the ferrite grain boundaries, and the material formability falls, so Mo is made 0.50% or less. Preferably it is 0.40% or less.
B: 0.0004 to 0.01%
B is an element raising the hardenability and further raising the toughness. In the steel sheet of the present invention, a predetermined hardenability is required, so 0.0004 to 0.01% is added. If less than 0.0004%, the effect of addition is not obtained, so B is made 0.0004% or more. Preferably it is 0.0010% or more.
On the other hand, if over 0.01%, coarse B compounds becoming the cause of internal defects and other flaws at the time of steel production are formed, so B is made 0.01% or less. Preferably it is 0.007% or less.
Ti: 0.001 to 0.10%
Ti is an element forming nitrides and contributing to refinement of the crystal grains and works to effectively bring out the effect of addition of B. If less than 0.001%, the effect of addition is not obtained, so Ti is made 0.001% or more. Preferably it is 0.010% or more.
On the other hand, if over 0.10%, coarse Ti nitrides are formed and the material formability falls, so Ti is made 0.10% or less. Preferably it is 0.07% or less.
The following elements are impurities and have to be controlled to certain amounts or less.
P: 0.02% or less
P is an element segregating at the ferrite grain boundaries and working to suppress the formation of carbides at the ferrite grain boundaries. For this reason, the smaller the amount of P, the better. The content of P may also be 0, but if reducing it to less than 0.0001%, the refining costs greatly increase, so the substantive lower limit is 0.0001 to 0.0013%.
If P is over 0.02%, formation of carbides at the ferrite grain boundaries is suppressed, the number of carbides decreases, and the material formability falls, so P is made 0.02% or less. Preferably it is 0.01% or less.
S: 0.01% or less
S is an impurity element forming MnS and other nonmetallic inclusions. The nonmetallic inclusions form starting points of fracture at the time of cold forging, so the smaller the S, the better. The content of S may also be 0, but to lower S to less than 0.0001%, the refining costs greatly increase, so the substantive lower limit is 0.0001 to 0.0012%.
If S is over 0.01%, nonmetallic inclusions are formed and the material formability falls, so S is made 0.01% or less. Preferably it is 0.009% or less.
N: 0.02% or less
N is an element which, if present in a large amount, causes embrittlement of the ferrite. For this reason, the smaller the amount of N, the better. The content of N may also be 0, but to lower N to less than 0.0001%, the refining costs greatly increase, so the substantive lower limit is 0.0001 to 0.0006%.
If N is over 0.02%, the ferrite becomes brittle and the material formability falls, so N is made 0.02% or less. Preferably it is 0.017% or less.
When the steel sheet of the present invention contains C: 0.10 to 0.40% and Mn: 0.30 to 1.00%, embrittlement of the ferrite is suppressed, so N is made 0.01% or less. Preferably it is 0.007% or less.
O: 0.02% or less O is an element which, if present in a large amount, promotes the formation of coarse oxides. For this reason, the smaller the amount of O, the better, but to lower O to less than 0.0001%, the refining costs greatly increase, so the amount is made 0.0001% or more. Preferably it is 0.0011% or more.
On the other hand, if over 0.020%, coarse oxides are formed in the steel, the oxides become starting points of fracture at the time of cold forging, and the material formability falls, so O is made 0.02% or less. Preferably it is 0.01% or less.
Sn: 0.05% or less
Sn is an element which unavoidably enters from the steel starting materials. For this reason, the smaller the amount of Sn, the better. The content of S may also be 0, but to lower S to less than 0.001%, the refining costs greatly increase, so the substantive lower limit is 0.001 to 0.002%.
On the other hand, if over 0.05%, the ferrite becomes brittle and the material formability falls, so Sn is made 0.05% or less. Preferably it is 0.04% or less.
Sb: 0.05% or less
Sb, like Sn, is an element which unavoidably enters from the steel starting materials, segregates at the ferrite grain boundaries, and reduces the number of carbides at the ferrite grain boundaries. For this reason, the smaller the amount of Sb, the better. The content of Sb may also be 0, but to lower Sb to less than 0.001%, the refining costs greatly increase, so the substantive lower limit is 0.001 to 0.002%.
On the other hand, if over 0.050%, Sb segregates at the ferrite grain boundaries, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so Sb is made 0.050% or less. Preferably it is 0.04% or less.
As: 0.05% or less
As, like Sn and Sb, is an element which unavoidably enters from the steel starting materials and segregates at the ferrite grain boundaries. For this reason, the smaller the amount of As, the better. The content of As may also be 0, but to lower As to less than 0.001%, the refining costs greatly increase, so the substantive lower limit is 0.001 to 0.002%.
On the other hand, if over 0.05%, As segregates at the ferrite grain boundaries, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so As is made 50% or less. Preferably it is 0.04% or less.
The steel sheet of the present invention has the above elements as basic elements, but may further contain the following elements for the purpose of improving the cold forgeability of the steel sheet. The following elements are not essential for obtaining the effects of the present invention, so the contents may also be 0.
Nb: 0.10% or less
Nb is an element effective for control of the form of the carbides. Further, it is an element refining the structure and contributing to improvement of the toughness. To obtain this effect of addition, Nb preferably is made 0.001% or more. More preferably it is 0.002% or more.
On the other hand, if over 0.10%, a large number of fine Nb carbides precipitate, the strength excessively rises, and, further, the number of carbides at the grain boundaries falls and the cold forgeability falls, so Nb is made 0.10% or less. Preferably it is 0.09% or less.
V: 0.10% or less
V, like Nb, is an element effective for control of the form of the carbides. Further, it is an element refining the structure and contributing to improvement of the toughness. To obtain this effect of addition, V preferably is made 0.01% or more. More preferably it is 0.004% or more.
On the other hand, if over 0.10%, a large number of fine V carbides are formed, the strength rises too much, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so V is made 0.10% or less. Preferably it is 0.09% or less.
Cu: 0.10% or less
Cu is an element segregating at the ferrite grain boundaries. Further, it is an element forming fine precipitates and contributing to the improvement of strength. To obtain the effect of improvement of strength, Cu preferably is made 0.001% or more. More preferably it is 0.008% or more.
On the other hand, if over 0.10%, segregation at the ferrite grain boundaries invites red shortness and causes the productivity in hot rolling to fall, 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 control of the form of carbides. To obtain this effect of addition,
W preferably is made 0.001% or more. More preferably it is 0.003% or more.
On the other hand, if over 0.10%, a large number of fine W carbides are formed, the strength rises too much, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so W is made 0.10% or less. Preferably it is 0.08% or less.
Ta: 0.001 to 0.10%
Ta, like Nb, V, and W, is an element effective for control of the form of carbides. To obtain this effect of addition, Ta preferably is made 0.001% or more. More preferably it is 0.007% or more.
On the other hand, if over 0.10%, a large number of fine Ta carbides are formed, the strength rises too much, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so T is made 0.100% or less. Preferably it is 0.09% or less.
Ni: 0.10% or less
Ni is an element effective for improvement of the impact resistance characteristic of the formed part. To obtain this effect of addition, Ni preferably is made 0.001% or more. More preferably it is 0.002% or more.
On the other hand, if over 0.10%, the number of carbides at the ferrite grain boundaries decreases and the material formability falls, so Ni is made 0.10% or less. Preferably it is 0.09% or less.
Mg: 0.05% or less
Mg is an element which can control the form of sulfides by addition in a trace amount. To obtain this effect of addition, Mg preferably is made 0.0001% or more. More preferably it is 0.0008% or more.
On the other hand, if over 0.05%, the ferrite becomes brittle and the material formability falls, so Mg is made 0.05% or less. Preferably it is 0.04% or less.
Ca: 0.05% or less
Ca, like Mg, is an element which can control the form of sulfides by addition in a trace amount. To obtain this effect of addition, Ca preferably is made 0.001% or more. More preferably it is 0.003% or more.
On the other hand, if over 0.05%, coarse Ca oxides are formed and become starting points of fracture at the time of forming by cold forging, that is, the material formability falls, so Ca is made 0.05% or less. Preferably it is 0.04% or less.
Y: 0.05% or less
Y, like Mg and Ca, is an element which can control the form of sulfides by addition in a trace amount. To obtain this effect of addition, Y preferably is made 0.001% or more. More preferably it is 0.003% or more.
On the other hand, if over 0.05%, coarse Y oxides are formed and become starting points of fracture at the time of forming by cold forging, that is, the material formability falls, so Y is made 0.05% or less. Preferably it is 0.03% or less.
Zr: 0.05% or less
Zr, like Mg, Ca, and Y, is an element which can control the form of sulfides by addition in a trace amount. To obtain this effect of addition, Zr preferably is made 0.001% or more. More preferably it is 0.004% or more.
On the other hand, if over 0.05%, coarse Zr oxides are formed and become starting points of fracture at the time of forming by cold forging, that is, the material formability falls, so Zr is made 0.05% or less. Preferably it is 0.04% or less.
La: 0.05% or less
La is an element able to control the form of the sulfides by addition in a trace amount, but is an element which segregates at the grain boundaries and reduces the number of carbides at the ferrite grain boundaries. To obtain the effect of control of the form of sulfides, La preferably is made 0.001% or more. More preferably it is 0.003% or more.
On the other hand, if over 0.05%, La segregates at the ferrite grain boundaries, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so La is made 0.05% or less. Preferably it is 0.04% or less.
Ce: 0.05% or less
Ce, like La, is an element able to control the form of the sulfides by addition in a trace amount, but is an element which segregates at the grain boundaries and reduces the number of carbides at the ferrite grain boundaries. To obtain the effect of control of the form of sulfides, Ce preferably is made 0.001% or more. More preferably it is 0.003% or more.
On the other hand, if over 0.05%, Ce segregates at the ferrite grain boundaries, the number of carbides at the ferrite grain boundaries decreases, and the material formability falls, so Ce is made 0.050% or less. Preferably it is 0.04% or less.
The balance of the chemical composition is Fe and unavoidable impurities.
Next, the structure of the steel sheet of the present invention will be explained.
The structure of the steel sheet of the present invention is a structure substantially comprised of ferrite and carbides. Carbides are compounds of iron and carbon of cementite (Fe₃C) plus compounds of cementite in which Fe atoms are substituted by Mn, Cr, and other alloy elements and alloy carbides (M₂₃C₆, M₆C, MC, etc. [M: Fe and other metal elements added as alloys]).
When forming steel sheet into a predetermined part shape, a shear zone is formed at the macrostructure of the steel sheet and slip deformation occurs concentrated near the shear zone. In slip deformation, along with proliferation of dislocations, a region of a high dislocation density is formed near the shear zone. Along with the increase in the amount of strain imparted to the steel sheet, slip deformation is promoted and the dislocation density increases.
In cold forging, strong working is performed with an equivalent strain exceeding 1. For this reason, in conventional steel sheet, it was not possible to prevent the formation of voids and/or cracks along with the increase in dislocation density and was difficult to improve the cold forgeability. To solve this problem, it is effective to suppress the formation of a shear zone at the time of forming.
From the viewpoint of the microstructure, formation of a shear zone can be understood as the phenomenon of slip occurring at a certain one grain crossing the crystal grain boundary and being continuously propagated to the adjoining grain. Accordingly, to suppress the formation of a shear zone, it is necessary to prevent propagation of slip crossing crystal grain boundaries.
The carbides in steel sheet are strong particles inhibiting slip. By forming carbides at the ferrite grain boundaries, it becomes possible to prevent the propagation of slip crossing crystal grain boundaries and suppress the formation of a shear zone and improve the cold forgeability. Simultaneously, the steel sheet is also improved in formability.
The formability of steel sheet is largely due to the accumulation of strain inside the crystal grains (accumulation of dislocations). If propagation of strain to the adjoining crystal grains is blocked at the crystal grain boundaries, the amount of strain inside the crystal grains increases. As a result, the work hardening rate increases and the formability is improved.
To obtain such an effect, carbides have to be made to disperse in the metal structure in suitable sizes. Therefore, the average grain size of carbides is made 0.4 μm to 2.0 μm. If the average grain size of the carbides is less than 0.4 μm, the steel sheet remarkably increases in hardness and falls in cold forgeability.
More preferably it is 0.6 μm or more.
On the other hand, if the average particle size of the carbides exceeds 2.0 μm, at the time of cold forming, the carbides form starting points of cracks. More preferably, it is 1.95 μm or less.
Further, cementite, a carbide of iron, is a hard and brittle structure. If present in the form of pearlite, which is a layered structure with ferrite, the steel becomes hard and brittle. Therefore, pearlite has to be reduced as much as possible. In the steel sheet of the present invention, the area ratio is made 6% or less.
Pearlite has a unique lamellar structure, so can be discerned by observation by an SEM or optical microscope. By calculating the region of the lamellar structure at any cross-section, the area ratio of the pearlite can be found.
Based on theory and principle, cold forgeability is considered to be strongly affected by the rate of coverage of the ferrite grain boundaries by carbides. High precision measurement is sought, but measurement of the rate of coverage of ferrite grain boundaries by carbides in a three-dimensional space requires serial sectioning SEM observation using an FIB to repeatedly cut a sample for observation in a scanning electron microscope or 3D EBSP observation. A massive measurement time is required and technical knowhow has to be built up.
The inventors judged that the above method of observation was not a general method of analysis and did not employ it. They searched for a simpler, higher precision indicator for evaluation. As a result, they discovered that it is possible to quantitatively evaluate the cold forgeability and formability by using the ratio B/A of the number B of carbides at the ferrite grain boundaries to the number A of carbides in the ferrite grains as an indicator and that if the ratio B/A exceeds 1, the cold forgeability and the formability in drawing and thickening remarkably rise.
Buckling, folding, and twisting of the steel sheet occurring at the time of cold forging occur due to localization of strain accompanying the formation of a shear zone, so by forming carbides at the ferrite grain boundaries, the formation of a shear zone and localization of strain are reduced and buckling, folding, and twisting are suppressed.
The carbides are observed by a scanning electron microscope. Before observation, the sample for observation of the structure is polished by wet polishing by Emery paper and a diamond abrasive having an average particle size of 1 μm, the observed surface is polished to a mirror finish, then a 3% nitric acid-alcohol solution is used to etch the structure. The magnification of the observation was made 3000× and images of eight fields of 30 μm×40 μm at a sheet thickness ¼ layer were captured at random.
The obtained structural images were analyzed by image analyzing software (Win ROOF made by Mitani Shoji) to measure in detail the areas of the carbides contained in the analyzed regions. The circle equivalent diameters (=2×√(area/3.14)) were found from the areas of the carbides and the average value was made the grain size of the carbides. Note that, to keep down the effect of measurement error due to noise, carbides with an area of 0.01 μm²or less are excluded from the coverage of the evaluation.
The number of carbides present at the ferrite grain boundaries are counted, the number of carbides at the grain boundaries are subtracted from the total number of carbides, and the number of carbides in the ferrite grains are found. Based on the measured and calculated number of carbides, the ratio B/A of the number B of carbides at the ferrite grain boundaries with respect to the number A of carbides inside the ferrite grains is calculated.
In the structure of the steel sheet after annealing, the ferrite grain size is preferably 3 μm to 50 μm from the viewpoint of improvement of the cold forgeability. If the ferrite grain size is less than 3 μm, the hardness increases and fractures and cracks easily form at the time of cold forging, so the ferrite grain size is preferably 3 μm or more. More preferably it is 5 μm or more.
If the ferrite grain size is over 50.0 μm, the number of carbides on the crystal grain boundaries suppressing the propagation of slip is decreased and the cold forgeability falls, so the ferrite grain size is preferably 50 μm or less. More preferably it is 40 μm or less.
The ferrite grain size is measured by using the above-mentioned procedure to polish the observed surface of the sample surface to a mirror finish, then etching it by a 3% nitric acid-alcohol solution and observing the structure by an optical microscope or scanning electron microscope and applying the line segment method to the captured image.
At the time of cold forging, in addition to control of the form of carbides, drawability at the time of cold forging becomes necessary.
To improve the drawability at the time of cold forging, improvement of the plastic anisotropy becomes necessary. For this reason, control of the texture at the hot rolled steel sheet is necessary. The texture is evaluated by X-ray diffraction at a plane parallel to the sheet surface at a ½ sheet thickness part of the hot rolled steel sheet. For X-ray diffraction, X-rays from a Mo tube are used.
The diffraction intensities at the diffraction orientations {110}, {220}, {211}, and {310} due to reflection are obtained and based on these an ODF is prepared. For the preparation of the ODF, the diffraction intensity data of random orientations of iron is used. From this, the X-ray diffraction intensity of {211}<011>is found as I1 and the X-ray diffraction intensity of {100}<011>is found as I0. If this I1/I0 is less than 1, it means that the recrystallization necessary for a random texture appears at the time of hot rolling. If the random texture can be obtained, the plastic anisotropy is reduced and the formability is improved.
By making the Vickers hardness of the steel sheet 100 HV to 150 HV (when C: 0.10 to 0.40% and Mn: 0.01 to 0.30%) or by making it 100 HV to 170 HV, it is possible to improve the formability at the time of cold forging. If the Vickers hardness is less than 100 HV, buckling easily occurs during the forming at the time of cold forging and the shaped part falls in precision, so the Vickers hardness is made 100 HV or more. Preferably, it is 110 HV or more.
If the Vickers hardness is over 170 HV, the ductility falls, buckling to outside the plane easily occurs during thickening or other compression deformation, further, internal fracture easily occurs at the time of cold forging, and the impact resistance characteristic deteriorates, so the Vickers hardness is made 170 HV or less. To reliably secure the ductility and impact resistance characteristic, the Vickers hardness is preferably made 150 HV or less. More preferably, it is 140 HV or less.
Next, the method of production of steel sheet of the present invention will be explained.
The method of production of the present invention has as its basic idea to use a steel slab of the above-mentioned chemical composition and integrally manage the hot rolling conditions and annealing conditions to control the structure of the steel sheet.
First, a steel slab obtained by continuously casting molten steel of the required chemical composition is used for hot rolling. The continuously cast slab may be directly used for hot rolling or may be used for hot rolling after cooling once, then heating.
If cooling once, then heating the steel slab for use for hot rolling, the heating temperature is preferably 1000° C. to 1250° C. and the heating time is preferably 0.5 hour to 3 hours. If directly using the continuously cast steel slab for hot rolling, the temperature of the steel slab used for the hot rolling is preferably made 1000° C. to 1250° C.
If the temperature of the steel slab or the heating temperature of the steel slab is over 1250° C. or the heating time of the steel slab is over 3 hours, decarburization from the surface layer of the steel slab becomes remarkable, at the time of heating before carburization and quenching, the austenite grains at the surface layer of the steel sheet abnormally grow, and the impact resistance falls. For this reason, the temperature of the steel slab or the heating temperature of the steel slab is preferably 1250° C. or less and the heating time is preferably 3 hours or less. More preferably, they are 1200° C. or less and 2.5 hours or less.
If the temperature of the steel slab or the heating temperature of the steel slab is less than 1000° C. or the heating time is less than 0.5 hour, the microsegregation and macrosegregation occurring in casting cannot be eliminated, regions remain inside the steel slab where Si, Mn, and other alloy elements locally concentrate, and the impact resistance falls. For this reason, the temperature of the steel slab or the heating temperature of the steel slab is preferably 1000° C. or more and the heating time is preferably 0.5 hour or more. More preferably they are 1050° C. or more and 1 hour or more.
The finish rolling in the hot rolling is completed at 820° C. or more, preferably at 900° C. to 950° C. in temperature region. If the finish rolling temperature is less than 820° C., the steel sheet increases in deformation resistance, the rolling load remarkably rises, and, further, the amount of roll wear increases and the productivity falls. Along with this, the recrystallization required for improving the plastic anisotropy does not sufficiently proceed, so the finish rolling temperature is made 820° C. or more. From the viewpoint of promoting recrystallization, it is preferably 900° C. or more.
If the finish rolling temperature is over 950° C., bulky scale forms during passage through the run out table (ROT). Due to this scale, flaws are formed at the surface of the steel sheet. When an impact load is applied after cold forging and carburization, quenching, and tempering, cracks easily form starting from the flaws, so the steel sheet falls in impact resistance. For this reason, the finish rolling temperature is made 950° C. or less. Preferably it is 920° C. or less.
When cooling the hot rolled steel sheet after finish rolling at the ROT, the cooling rate is preferably 10° C/sec to 100° C/sec. If the cooling rate is less than 10° C/sec, bulky scale is formed during the cooling. It is not possible to suppress the formation of flaws due to this and the impact resistance falls, so the cooling rate is preferably 10° C/sec or more. More preferably it is 15° C/sec or more.
If cooling from the surface layer of the steel sheet to the inside by an over 100° C/sec cooling rate, the outermost layer part is excessively heated and bainite, martensite, and other low temperature transformed structures are formed. When coiling, then cooling down to 100° C. to room temperature, then paying out the hot rolled steel sheet coil, microcracks form in the low temperature transformed structures. The microcracks are difficult to remove by pickling and cold rolling.
Further, if applying an impact load to the steel sheet after cold forging and carburization, quenching, and tempering, cracks advance starting from the microcracks, so the impact resistance falls. For this reason, to suppress the formation of bainite, martensite, and other low temperature transformed structures at the outermost layer part of the steel sheet, the cooling rate is preferably 100° C/sec or less. More preferably it is 90° C/sec or less.
Note that, the cooling rate indicates the cooling ability received from the cooling facilities in a water spray section at the time when being cooled on the ROT down to the target temperature of coiling from the time when the hot rolled steel sheet after finish rolling is water cooled at a water spray section after passing through a non-water spray section. It does not show the average cooling rate from the starting point of water spray to the temperature at which the steel sheet is coiled up by the coiler.
The coiling temperature is made 400° C. to 550° C. This is a temperature lower than the general coiling temperature and in particular is a condition not generally used when the content of C is high. By coiling up the hot rolled steel sheet produced under the above conditions in this temperature range, the structure of the steel sheet can be made a bainite structure comprised of fine ferrite in which carbides are dispersed.
If the coiling temperature is less than 400° C., the austenite, which was not transformed before coiling, transforms to hard martensite. At the time of paying out the hot rolled steel sheet coil, cracks form at the surface layer of the hot rolled steel sheet and the impact resistance falls.
Furthermore, at the time of recrystallization from austenite to ferrite, since the recrystallization driving force is small, the recrystallized ferrite grains are strongly influenced in orientation by the orientation of the austenite grains and randomization of the texture becomes difficult. For this reason, the coiling temperature is made 400° C. or more. Preferably it is 430° C. or more.
If the coiling temperature is over 550° C., pearlite with the large lamellar spacing is formed and highly heat stable bulky needle-shaped carbides are formed. These needle-shaped carbides remain even after two-stage annealing. At the time of cold forging and otherwise forming steel sheet, cracks are formed starting from these needle-shaped carbides.
Further, at the time of recrystallization from austenite to ferrite, conversely the recrystallization driving force becomes too large. In this case as well, the result becomes recrystallized ferrite grains heavily dependent on the orientation of the austenite grains and the texture is not randomized. For this reason, the coiling temperature is made 550° C. or less. Preferably it is 520° C. or less.
The hot rolled steel sheet coil is paid out and pickled, then is held in two temperature regions for two-stage step type of annealing (two-stage annealing). By treating the hot rolled steel sheet by two-stage annealing, the stability of the carbides is controlled to promote the formation of carbides at the ferrite grain boundaries.
If cold rolling the pickled steel sheet before annealing treatment, the ferrite grains are refined, so the steel sheet becomes harder to soften. For this reason, in the present invention, it is not preferable to cold roll the steel before annealing. It is preferable to perform the annealing treatment without cold rolling after the pickling.
The first stage of annealing is performed at 650 to 720° C., preferably the A_c1point or less in temperature region. Due to this annealing, the carbides are coarsened and partially spheroidized and the alloy elements are made to concentrate at the carbides to thereby raise the thermal stability of the carbides.
In the first stage of annealing, the heating rate up to the annealing temperature (below, referred to as the “first stage heating rate”) is made 30° C/hour to 150° C/hour. If the first stage heating rate is less than 30° C/hour, raising the temperature takes time and the productivity falls, so the first stage heating rate is made 3° C/hour or more. Preferably it is 10° C/hour or more.
On the other hand, if the first stage heating rate is over 150° C/hour, the temperature difference between the outer circumferential part and the inside part of the hot rolled steel sheet coil increases, scratches and seizing occur due to the difference in heat expansion, and relief shapes are formed at the steel sheet surface. At the time of cold forging and other forming, cracks occur starting from the relief shapes and result in a drop in cold forgeability and a drop in formability and impact resistance after carburizing, quenching, and tempering, so the first stage heating rate is made 150° C/hour or less. Preferably, it is 130° C/hour or less.
The annealing temperature in the first stage of annealing (below, referred to as “the first stage annealing temperature”) is made 650° C. to 720° C. If the first stage annealing temperature is less than 650° C., the carbides become insufficient in stability and it becomes difficult to form carbides remaining in the austenite in the second stage of annealing. Therefore, the first stage annealing temperature is made 650° C. or more. Preferably it is 670° C. or more.
On the other hand, if the annealing temperature exceeds 720° C., before the carbides rise in stability, austenite is formed and it becomes impossible to control the above-mentioned changes in structure, so the first stage annealing temperature is made 720° C. or less. Preferably it is 700° C. or less.
The annealing time in the first stage of annealing (below, referred to as the “first stage annealing time”) is made 3 hours to 60 hours. If the first stage annealing time is less than 3 hours, the carbides become insufficient in stability and it becomes difficult to form carbides remaining in the second stage of annealing. Therefore, the first stage annealing time is made 3 hours or more. Preferably it is 5 hours or more.
On the other hand, if the first stage annealing time exceeds 60 hours, no further stabilization of the carbides can be expected. Furthermore, the productivity drops. Therefore, the first stage annealing time is made 60 hours or less. Preferably it is 55 hours or less.
After that, the temperature is raised to 725 to 790° C., preferably the A_c1point to the A₃point in temperature region, to form austenite in the structure. At this time, the carbides in the fine ferrite grains dissolve in the austenite, but the carbides coarsened due to the first stage of annealing remain in the austenite.
When cooling without performing the second stage of annealing, the ferrite grain size does not become larger and the ideal structure cannot be obtained.
The heating rate up to the annealing temperature in the second stage of annealing (below, referred to as the “second stage heating rate”) is made 1° C/hour to 80° C/hour. At the time of the second stage of annealing, austenite is formed and grows from the ferrite grain boundaries. At that time, by slowing the heating rate up to the annealing temperature, it becomes possible to suppress the formation of nuclei of austenite raise the rate of coverage of the grain boundaries by carbides in the structure formed by gradual cooling after annealing.
For this reason, the second stage heating rate preferably is slow, but if less than 1° C/hour, raising the temperature takes time and the productivity falls, so the second stage heating rate is made 1° C/hour or more. Preferably it is 10° C/hour or more.
If the second stage heating rate is over 80° C/hour, the temperature difference between the outer circumferential part and the inside part of the hot rolled steel sheet coil increases, scratches and seizing occur due to the large difference in heat expansion due to the transformation, and relief shapes are formed at the steel sheet surface. At the time of cold forging, cracks occur starting from the relief shapes and result in a drop in cold forgeability and formability and further a drop in the impact resistance after carburizing, quenching, and tempering as well, so the second stage heating rate is made 80° C/hour or less. Preferably, it is 70° C/hour or less.
The annealing temperature in the second stage of annealing (below, referred to as “the second stage annealing temperature”) is made 725° C. to 790° C. If the second stage annealing temperature is less than 725° C., the amount of austenite formed becomes smaller, the number of carbides at the ferrite grain boundaries decreases after cooling after the second stage of annealing, and, further, the ferrite grain size becomes smaller. Therefore, the second stage annealing temperature is made 725° C. or more. Preferably it is 735° C. or more.
On the other hand, if the second stage annealing temperature exceeds 790° C., it becomes difficult to make carbides remain in the austenite and it becomes difficult to control the changes in structure, so the second stage annealing temperature is made 790° C. or less. Preferably it is 770° C. or less.
The annealing time in the second stage of annealing (“second stage annealing time”) is made 3 hours to 10 hours. If the second stage annealing time is less than 3 hours, the amount of formation of austenite becomes small, the carbides inside the ferrite grains do not sufficiently dissolve, it becomes difficult to make the number of carbides at the ferrite grain boundaries increase, and, further, the ferrite grain size becomes small. Therefore, the second stage annealing time is made 3 hours or more. Preferably it is 5 hours or more.
On the other hand, if the second stage annealing time exceeds 10 hours, it becomes difficult to make carbides remain in the austenite. Further, the manufacturing costs also increase. Therefore, the second stage annealing time is made less than 10 hours. Preferably it is 8 hours or less.
After the two-stage annealing, the steel sheet is cooled by a 1° C./ hour to 100° C/hour cooling rate down to 650° C.
By gradually cooling the austenite formed at the second stage of annealing due to gradual cooling, it transforms to ferrite, carbon atoms are adsorbed at the carbides remaining in the austenite, the carbides and austenite cover the ferrite grain boundaries, and, finally, a structure can be obtained in which a large amount of carbides are present at the ferrite grain boundaries. For this reason, the cooling rate is preferably slow, but if less than 1° C/hour, the time required for cooling increases and the productivity falls, so the cooling rate is made 1° C/hour or more.
Preferably it is 10° C/hour or more.
On the other hand, if the cooling rate is over 100° C/hour, the austenite transforms to pearlite, the steel sheet increases in hardness, and the cold forgeability falls. Further, after carburization, quenching, and tempering, the impact resistance falls. Therefore, the cooling rate is made 100° C/hour or less. Preferably it is 80° C/hour or less.
Furthermore, after cooling it down to 650° C., the steel sheet is cooled down to room temperature. The cooling rate at this time is not limited.
The atmosphere in the two-stage annealing is not particularly limited. For example, it may be any of a 95% or more nitrogen atmosphere, 95% or more hydrogen atmosphere, or air atmosphere.
As explained above, according to the method of production of the present invention integrally managing the hot rolling conditions and annealing conditions and controlling the structure of the steel sheet, it is possible to produce steel sheet excellent in formability at the time of cold forging combining drawing and thickening and, furthermore, excellent in the hardenability required for improvement of the impact resistance after carburization, quenching, and tempering. Examples
Next, examples of the present invention will be explained, but the conditions in the examples are illustrations of conditions employed for confirming the workability and effects of the present invention. The present invention is not limited to these illustrations of conditions. The present invention can employ various conditions so long as not departing from the gist of the present invention and achieving the object of the present invention.

EXAMPLE 1

A continuously cast slab (steel slab) of each of the chemical compositions shown in Table 1 and Table 2 (continuation of Table 1) was heated at 1240° C. for 1.8 hours, then hot rolled, cooled down to 530° C. by a 45° C/sec cooling rate on the ROT after finish hot rolling at 920° C., and coiled at 520° C. to produce sheet thickness 5.2 mm hot rolled steel sheet coil.
The hot rolled steel sheet coil was paid out and pickled, then was loaded into a box type annealing furnace. The annealing atmosphere was controlled to 95% hydrogen −5% nitrogen, then the coil was heated from room temperature to 705° C. by a 100° C/hour heating rate and was held at 710° C. for 24 hours to obtain a uniform temperature distribution inside the hot rolled steel sheet coil.
Next, the coil was heated by a 5° C/hour heating rate to 740° C., was further held at 740° C. for 5 hours, then was cooled down to 650° C. by a 10° C/hour cooling rate, then was furnace cooled down to room temperature to prepare a sample for evaluation of performance. The structure of the sample was observed by the method explained above and the ferrite grain size and number of carbides were measured.

TABLE 1

No	C	Si	Mn	P	S	Al	N	O	Ti	Cr	Mo	B	Nb	V	Cu	Remarks

1	0.35	0.21	0.53	0.0105	0.0027	0.065	0.0067	0.01	0.041	0.0300	0.1760	0.0090				Inv. steel
2	0.24	0.22	0.59	0.0075	0.0040	0.066	0.0021	0.01	0.035	0.4500	0.1330	0.0055				Inv. steel
3	0.18	0.03	0.41	0.0165	0.0026	0.029	0.0050	0.01	0.094	0.1500	0.4390	0.0022				Inv. steel
4	0.31	0.13	0.31	0.0003	0.0053	0.072	0.0048	0.02	0.050	0.1300	0.2510	0.0028				Inv. steel
5	0.25	0.09	0.39	0.0021	0.0088	0.055	0.0079	0.01	0.085	0.0500	0.3580	0.0098				Inv. steel
6	0.13	0.04	0.7	0.0088	0.0066	0.072	0.0031	0.02	0.058	0.4900	0.0170	0.0018				Inv. steel
7	0.14	0.04	0.32	0.0133	0.0081	0.052	0.0017	0.02	0.094	0.3700	0.3190	0.0093				Inv. steel
8	0.21	0.03	0.93	0.0147	0.0077	0.074	0.0021	0.01	0.085	0.3300	0.1440	0.0042				Inv. steel
9	0.29	0.22	0.4	0.0118	0.0082	0.036	0.0071	0.01	0.040	0.2700	0.0170	0.0082				Inv. steel
10	0.31	0.23	0.6	0.0040	0.0068	0.037	0.0010	0.01	0.039	0.2400	0.4480	0.0020				Inv. steel
11	1.20	0.02	0.78	0.0176	0.0049	0.073	0.0092	0.00	0.020	0.1600	0.0070	0.0024				Comp. steel
12	0.27	1.50	1.0	0.0170	0.0029	0.013	0.0004	0.00	0.037	0.2900	0.2990	0.0096				Comp. steel
13	0.37	0.21	3.3	0.0189	0.0052	0.046	0.0099	0.01	0.027	0.1400	0.2770	0.0077				Comp. steel
14	0.39	0.27	0.83	0.0091	0.0059	0.018	0.0023	0.01	0.038	1.1000	0.4070	0.0099				Comp. steel
15	0.35	0.26	0.37	0.0143	0.0031	0.018	0.0063	0.00	0.028	0.0200	0.1340	0.0000				Comp. steel
16	0.22	0.05	0.36	0.0109	0.0032	0.067	0.0008	0.0062	0.001	0.18	0.081	0.0014	0.028	0.083		Inv. steel
17	0.35	0.2	0.47	0.0137	0.0004	0.052	0.0013	0.0127	0.054	0.45	0.241	0.0057				Inv. steel
18	0.37	0.07	0.95	0.0076	0.0008	0.057	0.0098	0.0149	0.097	0.23	0.074	0.0019				Inv. steel
19	0.37	0.23	0.76	0.0086	0.0058	0.073	0.0057	0.0076	0.076	0.18	0.149	0.0007				Inv. steel
20	0.32	0.17	0.61	0.0197	0.0024	0.011	0.0009	0.0153	0.083	0.27	0.295	0.0019				Inv. steel
21	0.22	0.27	0.57	0.0059	0.0013	0.048	0.0041	0.0064	0.005	0.33	0.438	0.0085				Inv. steel
22	0.12	0.19	0.32	0.0135	0.0032	0.078	0.0096	0.0107	0.083	0.06	0.167	0.0051			0.024	Inv. steel
23	0.25	0.13	0.8	0.0056	0.0046	0.087	0.0064	0.0001	0.04	0.32	0.33	0.0079				Inv. steel
24	0.31	0.16	0.41	0.0142	0.0097	0.026	0.0054	0.0081	0.003	0.14	0.473	0.0049				Inv. steel
25	0.22	0.23	0.42	0.0052	0.003	0.093	0.0055	0.0197	0.05	0.21	0.349	0.0081				Inv. steel
26	0.34	0.24	0.68	0.01	0.0081	0.024	0.0077	0.0092	0.085	0.05	0.212	0.0024	0.044	0.09		Comp. steel
27	0.11	0.09	0.45	0.0124	0.0068	0.008	0.0059	0.0151	0.044	0.41	0.109	0.0074			0.21	Comp. steel
28	0.2	0.28	0.5	0.0196	0.0007	0.003	0.0060	0.0137	0.012	0.3	0.406	0.0013				Comp. steel
29	0.34	0.1	0.71	0.0004	0.003	0.014	0.0059	0.0093	0.092	0.02	0.189	0.0082		0.22	0.004	Comp. steel
30	0.37	0.09	0.8	0.007	0.0071	0.043	0.0032	0.0023	0.016	0.33	0.253	0.0044	0.36			Comp. steel
31	0.33	0.21	0.53	0.0104	0.0028	0.066	0.0069	0.01	0.039	0.0320	0.1730	0.0093				Inv. steel
32	0.36	0.04	0.32	0.011	0.0032	0.067	0.0008	0.0062	0.001	0.18	0.081	0.0014	0.054	0.087		Inv. steel
33	0.34	0.20	1.21	0.0105	0.0027	0.065	0.0075	0.01	0.041	0.0320	0.1720	0.0095				Inv. steel
34	0.22	0.05	1.83	0.01	0.0032	0.066	0.0007	0.0062	0.001	0.18	0.081	0.0014	0.032	0.096		Inv. steel
35	0.37	0.21	1.3	0.0104	0.0029	0.065	0.0063	0.01	0.038	0.0280	0.1690	0.0096				Inv. steel
36	0.22	0.06	2.19	0.0109	0.0035	0.067	0.0008	0.0062	0.001	0.18	0.081	0.0014	0.036	0.065		Inv. steel

TABLE 2

(Continuation of Table 1) (mass %)

No

W

Ta

Ni

Sn

Sb

As

Mg

Ca

Y

Zr

La

Ce

Remarks

1													Inv. steel
2													Inv. steel
3													Inv. steel
4													Inv. steel
5													Inv. steel
6													Inv. steel
7													Inv. steel
8													Inv. steel
9													Inv. steel
10													Inv. steel
11													Comp. steel
12													Comp. steel
13													Comp. steel
14													Comp. steel
15													Comp. steel
16													Inv. steel
17		0.033		0.043									Inv. steel
18	0.091												Inv. steel
19					0.014								Inv. steel
20						0.032			0.022				Inv. steel
21							0.0424						Inv. steel
22						0.031						0.046	Inv. steel
23											0.046		Inv. steel
24			0.017							0.024			Inv. steel
25								0.027					Inv. steel
26			0.15										Comp. steel
27		0.049						0.05					Comp. steel
28	0.33			0.012						0.041			Comp. steel
29									0.024				Comp. steel
30											0.002		Comp. steel
31													Inv. steel
32													Inv. steel
33													Inv. steel
34													Inv. steel
35													Inv. steel
36													Inv. steel

Table 3 shows the ferrite grain size (μm), average carbide grain size (μm), pearlite area ratio (%), Vickers hardness (HV), number of grain boundary carbides/number of grain carbides, X-ray intensity ratio I1/I0, “r” value anisotropy index |Δr|, and critical cooling rate (° C/sec) shown in Table 1 and Table 2. If I1/I0 is 1 or more, the recrystallization in hot rolling does not sufficiently proceed and the steel sheet becomes larger in plastic anisotropy. Note that, the “r” value anisotropy index |Δr| was found by a tensile test

TABLE 3

		Average			No. of grain
	Ferrite	carbide	Pearlite		boundary			Critical
	grain	area	grain	Vickers	carbides/No.			cooling
	size	size	rate	hardness	of grain			rate
No	(μm)	(μm)	(%)	(HV)	carbides	I1/I0	\|Δr\|	(° C./sec)	Remarks

1	17	0.7	1.7	120	5.66	0.75	0.17	29.8	Inv. ex.
2	23	1.1	1.2	118	5.85	0.81	0.19	29.8	Inv. ex.
3	16	0.8	0.8	105	3.88	0.80	0.18	29.9	Inv. ex.
4	13	1.1	1.3	109	5.54	0.76	0.17	30.0	Inv. ex.
5	20	0.9	1.0	108	3.48	0.69	0.15	29.9	Inv. ex.
6	12	1.1	0.0	115	7.11	0.67	0.14	29.9	Inv. ex.
7	13	1.3	1.0	100	6.11	0.64	0.13	30.1	Inv. ex.
8	11	1.2	1.3	124	4.23	0.64	0.13	29.7	Inv. ex.
9	20	1.1	1.2	113	6.07	0.79	0.18	29.9	Inv. ex.
10	17	1.4	1.9	122	6.59	0.65	0.13	29.8	Inv. ex.
11	18	1.1	9.1	167	5.36	0.69	0.15	30.0	Comp. ex.
12	10	1.0	1.5	154	6.86	0.69	0.15	29.4	Comp. ex.
13	14	1.3	12.3	178	5.69	0.72	0.16	13.2	Comp. ex.
14	20	1.0	1.2	133	0.91	0.78	0.18	29.6	Comp. ex.
15	12	1.1	1.9	114	6.04	0.79	0.18	311.0	Comp. ex.
16	11	0.8	1.8	105	5.02	0.63	0.13	30.1	Inv. ex.
17	26	1.0	1.6	119	4.19	0.66	0.14	29.8	Inv. ex.
18	13	1.2	1.0	134	5.05	0.78	0.18	29.6	Inv. ex.
19	11	1.2	1.8	130	5.36	0.75	0.17	29.6	Inv. ex.
20	20	1.1	1.0	122	4.17	0.65	0.13	29.7	Inv. ex.
21	11	1.1	1.0	116	5.57	0.74	0.16	29.9	Inv. ex.
22	17	1.1	0.6	100	6.17	0.82	0.19	30.1	Inv. ex.
23	14	1.2	1.0	122	6.19	0.77	0.17	29.7	Inv. ex.
24	12	1.1	1.8	113	7.12	0.81	0.19	29.9	Inv. ex.
25	17	0.8	1.4	111	5.38	0.74	0.16	29.9	Inv. ex.
26	13	1.2	1.4	126	6.46	0.80	0.18	29.6	Comp. ex.
27	19	1.1	1.2	105	6.93	0.74	0.16	30.0	Comp. ex.
28	23	0.9	1.8	113	6.58	0.70	0.15	29.9	Comp. ex.
29	18	1.2	0.5	125	4.62	0.71	0.15	29.7	Comp. ex.
30	14	1.1	0.7	129	4.93	0.68	0.14	29.7	Comp. ex.
31	18	1.1	1.5	125	5.72	0.75	0.17	29.3	Inv. ex.
32	13	0.8	0.2	109	5.17	0.63	0.13	29.7	Inv. ex.
33	18	1.3	1.2	125	5.60	0.75	0.16	31.2	Inv. ex.
34	13	1.4	1.9	109	4.97	0.63	0.13	30.9	Inv. ex.
35	18	1.3	1.4	128	5.75	0.75	0.17	29.5	Inv. ex.
36	13	1.4	1.0	110	5.23	0.63	0.11	29.8	Inv. ex.

In general, if the anisotropy index |Δr| obtained from the “r” values in parallel to the sheet surface and in three directions is over 0.2, the drawability falls. Therefore, to secure excellent formability, a |Δr| not over 2 is demanded.
The critical cooling rate was found by preparing a CCT graph. If cooling hot rolled steel sheet by a cooling rate slower than the found critical cooling rate, the hardenability at the time of hardening after forming a part becomes poorer and pearlite structures are formed so sufficient strength cannot be obtained. For this reason, the critical cooling rate must be small in order to obtain a high hardening strength. If the critical cooling rate is 280° C/sec, it can be judged that the hardenability is improved.
In the invention examples shown in Table 3, the average carbide grain size is 0.4 to 2.0 μm, the pearlite area ratio is 6% or less, the number of grain boundary carbides/number of grain carbides is over 1, and the I1/I0 is less than 1, so the Vickers hardness is 100 HV to 170 HV in range and |Δr| is less than 0.2. In the comparative examples using the comparative steel sheets, the Vickers hardness is over 150, while the number of grain boundary carbides/number of grain carbides becomes less than 1. In the comparative steel sheet in which B is not added (in Tables 1 and 2, No. 15), the critical cooling rate is over 280° C/sec and the hardenability falls.

EXAMPLE 2

A method of production of conditions outside the scope of conditions prescribed in the present invention was applied to the 12 types of steel of the Invention Steel Nos. 1 to 5, Nos. 16 to 19, Nos. 31, No. 33, and No. 35. Table 4 shows the manufacturing conditions, while Table 5 shows the ferrite grain size (μm), Vickers hardness (HV), number of grain boundary carbides/number of grain carbides, X-ray intensity ratio I1/I0, “r” value anisotropy index |Δr|, and critical cooling rate (° C/sec) of steel sheets produced under the manufacturing conditions shown in Table 4.

TABLE 4

	Hot rolling
	conditions	Annealing conditions

Finish

1st stage

2nd stage

Cooling

	rolling	Coiling	Heating	Holding	Holding	Heating	Holding	Holding	rate down
	temp.	temp.	rate	temp.	time	rate	temp.	time	to 650° C.
No	(° C.)	(° C.)	(° C./hour)	(° C.)	(hours)	(° C./hour)	(° C.)	(hours)	(° C./hour)

1	720	520	60	710	24	60	740	5	80
2	970	520	60	710	24	60	740	5	80
3	920	350	60	710	24	60	740	5	80
4	920	570	60	710	24	60	740	5	80
5	920	520	60	600	24	60	740	5	80
16	920	520	60	710	2	60	740	5	80
17	920	520	60	710	24	60	720	5	80
18	920	520	60	710	24	60	820	5	80
19	920	520	60	710	24	60	740	18	80
31	720	520	60	710	24	60	740	5	80
33	920	520	60	600	24	60	740	5	80
35	920	520	60	710	24	60	820	5	80

TABLE 5

		Average			No. of grain
	Ferrite	carbide	Pearlite		boundary			Critical
	grain	grain	area	Vickers	carbides/No.			cooling
	size	size	rate	hardness	of grain			rate
No	(μm)	(μm)	(%)	(HV)	carbides	I1/I0	\|Δr\|	(° C./sec)	Remarks

1	17	0.9	1.3	120	5.66	1.2	0.33	29.8	Comp. ex.
2	23	1.1	2.1	118	5.85	1.6	0.47	29.8	Comp. ex.
3	16	0.8	0.8	105	3.88	1.1	0.29	29.9	Comp. ex.
4	13	1.1	3.1	109	5.54	1.3	0.36	30.0	Comp. ex.
5	20	0.9	4.5	108	0.81	0.82	0.19	29.9	Comp. ex.
16	11	1.1	3.8	105	0.53	0.79	0.18	30.1	Comp. ex.
17	26	1.0	1.2	119	0.92	0.73	0.16	29.8	Comp. ex.
18	13	2.4	9.4	161	0.67	0.68	0.14	29.6	Comp. ex.
19	11	1.9	10.2	155	0.83	0.62	0.12	29.6	Comp. ex.
31	17	0.9	0.4	125	5.72	1.2	0.33	29.8	Comp. ex.
33	18	1.3	4.8	130	0.92	0.75	0.17	31.2	Comp. ex.
35	18	1.7	8.2	163	0.83	0.75	0.17	29.5	Comp. ex.

It will be understood that making the finish rolling temperature in hot rolling or the coiling temperature a temperature outside of the scope of conditions prescribed in the present invention invites a drop in the recrystallization and has a large effect on the randomization of the texture and as a result causes the value of |Δr| to rise. Further, it will be understood that if making the annealing conditions outside the scope of conditions prescribed in the present invention, the number of grain boundary carbides/number of grain boundary carbides becomes 1 or less and the state of distribution of carbides greatly changes.

INDUSTRIAL APPLICABILITY

As explained above, according to the present invention, it is possible to provide steel sheet excellent in hardenability and formability as a material and a method of production of the same. The steel sheet of the present invention is suitable for forming a part by cold forging such as thickening to obtain a gear or other part. Accordingly, the present invention has high applicability in the manufacture of steel sheet and industries utilizing it.

Claims

1. A steel sheet consisting of, by mass %,

C: 0.10 to 0.70%,

Si: 0.01 to 0.30%,

Mn: 0.30 to 3.00%,

Al: 0.001 to 0.10%,

Cr: 0.010 to 0.50%,

Mo: 0.0010 to 0.50%,

B: 0.0004 to 0.01%,

Ti: 0.001 to 0.10%,

P: 0.02% or less,

S: 0.01% or less,

N: 0.0200% or less,

O: 0.0200% or less,

Sn: 0.05% or less,

Sb: 0.05% or less,

As: 0.05% 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: 0.05% or less,

Ca: 0.05% or less,

Y: 0.05% or less,

Zr: 0.05% or less,

La: 0.05% or less, and

Ce: 0.05% or less and

a balance of Fe and unavoidable impurities,

wherein a metal structure of the steel sheet includes

carbide having an average grain size of 0.4 μm to 2.0 μm,

perlite having an area ratio of 6% or less and

ferrite wherein a ratio of a number of the carbides at ferrite grain boundaries to a number of the carbides in ferrite grains of over 1; and

I1/I0<1 is satisfied when an X-ray diffraction intensity at {211}<011>at a plane of a part of ½ sheet thickness of the steel sheet is denoted by “I1” and an X-ray diffraction intensity at {100}<011>is denoted by “I0”,

the steel sheet having a Vickers hardness of 100 HV to 150 HV.

2. A method of production for producing steel sheet according to claim 1 comprising

hot rolling a steel slab of a chemical composition according to claim 1 with finish rolling temperature between 820° C. and 950° C., to obtain hot rolled steel sheet;

coiling the hot rolled steel sheet at 400° C. to 550° C.;

pickling the coiled hot rolled steel sheet;

heating the pickled hot rolled steel sheet to an annealing temperature of 650° C. to 720° C. by a heating rate of 30° C/hour to 150° C/hour and holding the steel sheet for 3 hours to 60 hours as a first stage of annealing; next,

heating the hot rolled steel sheet to an annealing temperature of 725° C. to 790° C. by a heating rate of 1° C/hour to 80° C/hour and holding the steel sheet for 3 hours to less than 10 hours as a second stage of annealing; and, next,

cooling the annealed hot rolled steel sheet to 650° C. by a cooling rate of 1° C/hour to 100° C/hour.