U.S. patent application number 15/576682 was filed with the patent office on 2018-05-17 for steel sheet and method of production of same.
This patent application is currently assigned to NIPPON STEEL & SUMITOMO METAL CORPORATION. The applicant listed for this patent is NIPPON STEEL & SUMITOMO METAL CORPORATION. Invention is credited to Motonori HASHIMOTO, Kazuo HIKIDA, Ken TAKATA, Kengo TAKEDA.
Application Number | 20180135146 15/576682 |
Document ID | / |
Family ID | 57392862 |
Filed Date | 2018-05-17 |
United States Patent
Application |
20180135146 |
Kind Code |
A1 |
TAKATA; Ken ; et
al. |
May 17, 2018 |
STEEL SHEET AND METHOD OF PRODUCTION OF SAME
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 .mu.m to 2.0 .mu.m, 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 1/2 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.
Inventors: |
TAKATA; Ken; (Tokyo, JP)
; HIKIDA; Kazuo; (Tokyo, JP) ; TAKEDA; Kengo;
(Tokyo, JP) ; HASHIMOTO; Motonori; (Tokyo,
JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
NIPPON STEEL & SUMITOMO METAL CORPORATION |
Tokyo |
|
JP |
|
|
Assignee: |
NIPPON STEEL & SUMITOMO METAL
CORPORATION
Tokyo
JP
|
Family ID: |
57392862 |
Appl. No.: |
15/576682 |
Filed: |
May 26, 2016 |
PCT Filed: |
May 26, 2016 |
PCT NO: |
PCT/JP2016/065629 |
371 Date: |
November 22, 2017 |
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C22C 38/00 20130101;
C22C 38/32 20130101; C21D 6/005 20130101; C22C 38/28 20130101; C21D
6/008 20130101; C21D 8/0205 20130101; C21D 8/0226 20130101; C22C
38/06 20130101; C22C 38/24 20130101; C22C 38/60 20130101; C22C
38/002 20130101; C22C 38/02 20130101; C21D 8/02 20130101; C22C
38/20 20130101; C21D 2211/005 20130101; C22C 38/001 20130101; C23G
1/00 20130101; C21D 8/0263 20130101; C21D 2211/009 20130101; C21D
6/002 20130101; C21D 9/46 20130101; C22C 38/22 20130101; C22C 38/26
20130101; C22C 38/38 20130101 |
International
Class: |
C21D 9/46 20060101
C21D009/46; C22C 38/38 20060101 C22C038/38; C22C 38/32 20060101
C22C038/32; C22C 38/28 20060101 C22C038/28; C22C 38/26 20060101
C22C038/26; C22C 38/24 20060101 C22C038/24; C22C 38/22 20060101
C22C038/22; C22C 38/20 20060101 C22C038/20; C22C 38/06 20060101
C22C038/06; C22C 38/02 20060101 C22C038/02; C22C 38/00 20060101
C22C038/00; C21D 8/02 20060101 C21D008/02; C21D 6/00 20060101
C21D006/00 |
Foreign Application Data
Date |
Code |
Application Number |
May 26, 2015 |
JP |
2015-106755 |
Sep 28, 2015 |
JP |
2015-189883 |
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 .mu.m to 2.0 .mu.m,
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 1/2 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.degree. C. and 950.degree. C., to obtain hot rolled
steel sheet; coiling the hot rolled steel sheet at 400.degree. C.
to 550.degree. C.; pickling the coiled hot rolled steel sheet;
heating the pickled hot rolled steel sheet to an annealing
temperature of 650.degree. C. to 720.degree. C. by a heating rate
of 30.degree. C/hour to 150.degree. 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.degree. C. to 790.degree. C. by a heating rate of 1.degree.
C/hour to 80.degree. 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.degree. C. by a
cooling rate of 1.degree. C/hour to 100.degree. C/hour.
Description
TECHNICAL FIELD
[0001] The present invention relates to steel sheet and a method of
production of the same.
BACKGROUND ART
[0002] 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.
[0003] 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.
[0004] 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.
[0005] 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 .mu.m or more.
[0006] 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.
[0007] 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.
[0008] 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.
[0009] 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.
[0010] On the other hand, there have been the following proposals
for improvement of the formability, that is, the improvement of
plastic anisotropy.
[0011] 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
.DELTA.r.
[0012] 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 .DELTA.r 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.degree. C.
CITATION LIST
Patent Literature
[0013] PLT 1: Japanese Patent Publication No. 2013-040376A
[0014] PLT 2: Japanese Patent Publication No. 06-116679A
[0015] PLT 3: Japanese Patent Publication No. 09-201644A
[0016] PLT 4: Japanese Patent Publication No. 2006-213951A
[0017] PLT 5: Japanese Patent Publication No. 10-158780A
[0018] PLT 6: Japanese Patent Publication No. 2000-328172A
[0019] PLT 7: Japanese Patent Publication No. 2001-073076A
[0020] PLT 8: Japanese Patent Publication No. 2001-073077A
SUMMARY OF INVENTION
Technical Problem
[0021] 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.
[0022] 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
[0023] 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.
[0024] 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.
[0025] 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.
[0026] 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.
[0027] 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.
[0028] 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.degree. C. to
550.degree. 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.
[0029] 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.
[0030] 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.
[0031] The present invention was made based on these findings and
has as its gist the following:
[0032] (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,
[0033] 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 .mu.m to 2.0 .mu.m, 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 1/2 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.
[0034] (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.degree. C. and 950.degree. C., to obtain hot rolled
steel sheet; coiling the hot rolled steel sheet at 400.degree. C.
to 550.degree. C.; pickling the coiled hot rolled steel sheet;
heating the pickled hot rolled steel sheet to an annealing
temperature of 650.degree. C. to 720.degree. C. by a heating rate
of 30.degree. C/hour to 150.degree. 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.degree. C. to 790.degree. C. by a heating rate of 1.degree.
C/hour to 80.degree. 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.degree. C. by a
cooling rate of 1.degree. C/hour to 100.degree. C/hour.
Advantageous Effects of Invention
[0035] 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
[0036] 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 %".
[0037] C: 0.10 to 0.70%
[0038] 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.
[0039] 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.
[0040] 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.
[0041] Si: 0.01 to 0.30%
[0042] 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.
[0043] 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.
[0044] Mn: 0.30 to 3.00%
[0045] 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.
[0046] 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.
[0047] 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.
[0048] Al: 0.001 to 0.10%
[0049] 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.
[0050] 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.
[0051] Cr: 0.010 to 0.50%
[0052] 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.
[0053] 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.
[0054] Mo: 0.001 to 0.50%
[0055] 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.
[0056] 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.
[0057] B: 0.0004 to 0.01%
[0058] 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.
[0059] 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.
[0060] Ti: 0.001 to 0.10%
[0061] 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.
[0062] 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.
[0063] The following elements are impurities and have to be
controlled to certain amounts or less.
[0064] P: 0.02% or less
[0065] 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%.
[0066] 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.
[0067] S: 0.01% or less
[0068] 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%.
[0069] 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.
[0070] N: 0.02% or less
[0071] 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%.
[0072] 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.
[0073] 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.
[0074] 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.
[0075] 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.
[0076] Sn: 0.05% or less
[0077] 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%.
[0078] 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.
[0079] Sb: 0.05% or less
[0080] 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%.
[0081] 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.
[0082] As: 0.05% or less
[0083] 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%.
[0084] 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.
[0085] 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.
[0086] Nb: 0.10% or less
[0087] 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.
[0088] 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.
[0089] V: 0.10% or less
[0090] 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.
[0091] 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.
[0092] Cu: 0.10% or less
[0093] 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.
[0094] 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.
[0095] W: 0.10% or less
[0096] W, like Nb and V, is an element effective for control of the
form of carbides. To obtain this effect of addition,
[0097] W preferably is made 0.001% or more. More preferably it is
0.003% or more.
[0098] 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.
[0099] Ta: 0.001 to 0.10%
[0100] 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.
[0101] 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.
[0102] Ni: 0.10% or less
[0103] 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.
[0104] 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.
[0105] Mg: 0.05% or less
[0106] 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.
[0107] 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.
[0108] Ca: 0.05% or less
[0109] 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.
[0110] 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.
[0111] Y: 0.05% or less
[0112] 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.
[0113] 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.
[0114] Zr: 0.05% or less
[0115] 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.
[0116] 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.
[0117] La: 0.05% or less
[0118] 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.
[0119] 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.
[0120] Ce: 0.05% or less
[0121] 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.
[0122] 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.
[0123] The balance of the chemical composition is Fe and
unavoidable impurities.
[0124] Next, the structure of the steel sheet of the present
invention will be explained.
[0125] 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 (Fe.sub.3C)
plus compounds of cementite in which Fe atoms are substituted by
Mn, Cr, and other alloy elements and alloy carbides
(M.sub.23C.sub.6, M.sub.6C, MC, etc. [M: Fe and other metal
elements added as alloys]).
[0126] 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.
[0127] 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.
[0128] 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.
[0129] 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.
[0130] 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.
[0131] 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 .mu.m to 2.0 .mu.m. If
the average grain size of the carbides is less than 0.4 .mu.m, the
steel sheet remarkably increases in hardness and falls in cold
forgeability.
[0132] More preferably it is 0.6 .mu.m or more.
[0133] On the other hand, if the average particle size of the
carbides exceeds 2.0 .mu.m, at the time of cold forming, the
carbides form starting points of cracks. More preferably, it is
1.95 .mu.m or less.
[0134] 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.
[0135] 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.
[0136] 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.
[0137] 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.
[0138] 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.
[0139] 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 .mu.m, 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.times. and images of eight fields of 30
.mu.m.times.40 .mu.m at a sheet thickness 1/4 layer were captured
at random.
[0140] 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.times. (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
.mu.m.sup.2 or less are excluded from the coverage of the
evaluation.
[0141] 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.
[0142] In the structure of the steel sheet after annealing, the
ferrite grain size is preferably 3 .mu.m to 50 .mu.m from the
viewpoint of improvement of the cold forgeability. If the ferrite
grain size is less than 3 .mu.m, the hardness increases and
fractures and cracks easily form at the time of cold forging, so
the ferrite grain size is preferably 3 .mu.m or more. More
preferably it is 5 .mu.m or more.
[0143] If the ferrite grain size is over 50.0 .mu.m, 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 .mu.m or less. More
preferably it is 40 .mu.m or less.
[0144] 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.
[0145] At the time of cold forging, in addition to control of the
form of carbides, drawability at the time of cold forging becomes
necessary.
[0146] 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 1/2 sheet thickness part of the
hot rolled steel sheet. For X-ray diffraction, X-rays from a Mo
tube are used.
[0147] 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.
[0148] 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.
[0149] 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.
[0150] Next, the method of production of steel sheet of the present
invention will be explained.
[0151] 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.
[0152] 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.
[0153] If cooling once, then heating the steel slab for use for hot
rolling, the heating temperature is preferably 1000.degree. C. to
1250.degree. 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.degree. C. to 1250.degree. C.
[0154] If the temperature of the steel slab or the heating
temperature of the steel slab is over 1250.degree. 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.degree. C. or less and the heating time is
preferably 3 hours or less. More preferably, they are 1200.degree.
C. or less and 2.5 hours or less.
[0155] If the temperature of the steel slab or the heating
temperature of the steel slab is less than 1000.degree. 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.degree. C. or more
and the heating time is preferably 0.5 hour or more. More
preferably they are 1050.degree. C. or more and 1 hour or more.
[0156] The finish rolling in the hot rolling is completed at
820.degree. C. or more, preferably at 900.degree. C. to 950.degree.
C. in temperature region. If the finish rolling temperature is less
than 820.degree. 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.degree. C. or more. From the viewpoint of
promoting recrystallization, it is preferably 900.degree. C. or
more.
[0157] If the finish rolling temperature is over 950.degree. 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.degree. C. or less. Preferably it is 920.degree. C. or
less.
[0158] When cooling the hot rolled steel sheet after finish rolling
at the ROT, the cooling rate is preferably 10.degree. C/sec to
100.degree. C/sec. If the cooling rate is less than 10.degree.
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.degree.
C/sec or more. More preferably it is 15.degree. C/sec or more.
[0159] If cooling from the surface layer of the steel sheet to the
inside by an over 100.degree. 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.degree. 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.
[0160] 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.degree. C/sec or less. More preferably it is
90.degree. C/sec or less.
[0161] 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.
[0162] The coiling temperature is made 400.degree. C. to
550.degree. 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.
[0163] If the coiling temperature is less than 400.degree. 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.
[0164] 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.degree. C. or more. Preferably
it is 430.degree. C. or more.
[0165] If the coiling temperature is over 550.degree. 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.
[0166] 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.degree. C. or less.
Preferably it is 520.degree. C. or less.
[0167] 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.
[0168] 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.
[0169] The first stage of annealing is performed at 650 to
720.degree. C., preferably the A.sub.c1 point 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.
[0170] 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.degree. C/hour to 150.degree. C/hour. If
the first stage heating rate is less than 30.degree. C/hour,
raising the temperature takes time and the productivity falls, so
the first stage heating rate is made 3.degree. C/hour or more.
Preferably it is 10.degree. C/hour or more.
[0171] On the other hand, if the first stage heating rate is over
150.degree. 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.degree. C/hour or less.
Preferably, it is 130.degree. C/hour or less.
[0172] The annealing temperature in the first stage of annealing
(below, referred to as "the first stage annealing temperature") is
made 650.degree. C. to 720.degree. C. If the first stage annealing
temperature is less than 650.degree. 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.degree. C. or more. Preferably it is 670.degree. C. or
more.
[0173] On the other hand, if the annealing temperature exceeds
720.degree. 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.degree. C. or less. Preferably it is 700.degree. C. or
less.
[0174] 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.
[0175] 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.
[0176] After that, the temperature is raised to 725 to 790.degree.
C., preferably the A.sub.c1 point to the A.sub.3 point 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.
[0177] When cooling without performing the second stage of
annealing, the ferrite grain size does not become larger and the
ideal structure cannot be obtained.
[0178] 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.degree. C/hour to 80.degree. 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.
[0179] For this reason, the second stage heating rate preferably is
slow, but if less than 1.degree. C/hour, raising the temperature
takes time and the productivity falls, so the second stage heating
rate is made 1.degree. C/hour or more. Preferably it is 10.degree.
C/hour or more.
[0180] If the second stage heating rate is over 80.degree. 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.degree. C/hour or less.
Preferably, it is 70.degree. C/hour or less.
[0181] The annealing temperature in the second stage of annealing
(below, referred to as "the second stage annealing temperature") is
made 725.degree. C. to 790.degree. C. If the second stage annealing
temperature is less than 725.degree. 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.degree. C. or more. Preferably it is 735.degree. C. or
more.
[0182] On the other hand, if the second stage annealing temperature
exceeds 790.degree. 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.degree. C. or less. Preferably it is 770.degree. C. or
less.
[0183] 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.
[0184] 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.
[0185] After the two-stage annealing, the steel sheet is cooled by
a 1.degree. C./ hour to 100.degree. C/hour cooling rate down to
650.degree. C.
[0186] 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.degree. C/hour, the time required for cooling
increases and the productivity falls, so the cooling rate is made
1.degree. C/hour or more.
[0187] Preferably it is 10.degree. C/hour or more.
[0188] On the other hand, if the cooling rate is over 100.degree.
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.degree.
C/hour or less. Preferably it is 80.degree. C/hour or less.
[0189] Furthermore, after cooling it down to 650.degree. C., the
steel sheet is cooled down to room temperature. The cooling rate at
this time is not limited.
[0190] 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.
[0191] 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
[0192] 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
[0193] 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.degree. C. for 1.8 hours, then hot
rolled, cooled down to 530.degree. C. by a 45.degree. C/sec cooling
rate on the ROT after finish hot rolling at 920.degree. C., and
coiled at 520.degree. C. to produce sheet thickness 5.2 mm hot
rolled steel sheet coil.
[0194] 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.degree. C. by a
100.degree. C/hour heating rate and was held at 710.degree. C. for
24 hours to obtain a uniform temperature distribution inside the
hot rolled steel sheet coil.
[0195] Next, the coil was heated by a 5.degree. C/hour heating rate
to 740.degree. C., was further held at 740.degree. C. for 5 hours,
then was cooled down to 650.degree. C. by a 10.degree. 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-US-00001 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-US-00002 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
[0196] Table 3 shows the ferrite grain size (.mu.m), average
carbide grain size (.mu.m), 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
|.DELTA.r|, and critical cooling rate (.degree. 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 |.DELTA.r| was found by a tensile test
TABLE-US-00003 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 (.mu.m) (.mu.m)
(%) (HV) carbides I1/I0 |.DELTA.r| (.degree. 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.
[0197] In general, if the anisotropy index |.DELTA.r| 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 |.DELTA.r| not over 2 is demanded.
[0198] 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.degree. C/sec, it can be judged that the hardenability is
improved.
[0199] In the invention examples shown in Table 3, the average
carbide grain size is 0.4 to 2.0 .mu.m, 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 |.DELTA.r| 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.degree. C/sec
and the hardenability falls.
EXAMPLE 2
[0200] 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 (.mu.m),
Vickers hardness (HV), number of grain boundary carbides/number of
grain carbides, X-ray intensity ratio I1/I0, "r" value anisotropy
index |.DELTA.r|, and critical cooling rate (.degree. C/sec) of
steel sheets produced under the manufacturing conditions shown in
Table 4.
TABLE-US-00004 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.degree. C. No (.degree. C.) (.degree.
C.) (.degree. C./hour) (.degree. C.) (hours) (.degree. C./hour)
(.degree. C.) (hours) (.degree. 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-US-00005 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 (.mu.m) (.mu.m)
(%) (HV) carbides I1/I0 |.DELTA.r| (.degree. 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.
[0201] 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 |.DELTA.r| 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
[0202] 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.
* * * * *