The present invention relates to a steel sheet with excellent cold workability during
forming and a method for manufacturing the sheet.
Background Art
[0002]
Automotive parts, knives, and other mechanical parts are manufactured through
working processes sucb as punching, bending, and pressing. In the working processes,
improvement of workability is required for a material carbon steel sheet, in order to improve
product quality and stability and/or cost reduction.
[0003]
Generally, a carbon steel sheet is subjected to cold rolling and spheroidizing annealing,
so as to produce a soft carbon steel sheet with excellent workability made of ferrite and
spheroidized carbide. Many technologies for improving the workability of carbon steel sheets
have been proposed so far.
[0004]
For example, Patent Document 1 discloses a high-carbon steel sheet for precision
punching and a method for producing the sheet, wherein the sheet comprises, in terms of% by
mass, C: 0.15 to 0.90%, Si: 0.40% or less, Mn: 0.3 to LO%, P: 0.03% or less, total Al: 0.1% or
less, Ti: 0.01 to 0.05%, B: 0.0005 to 0.0050%, N: 0.01% or less, and Cr: 1.2% or less, has a
structure in which carbides having an average carbide grain size of0.4 to 1.0 f.tm and a carbide
spheroidization ratio of 80% or more are dispersed in a ferrite matrix, and has a notched
tensile elongation of20% or more.
[0005]
Patent Document 2 discloses a medium- to high-carbon steel sheet with excellent
workability and a method for producing the sheet, wherein the sheet comprises C: 0.3 to 1.3
wt%, Si: 1.0 wt% or less, Mn: 0.2 to 1.5 wt%, P: 0.02 wt% or less, and S: 0.02 wt% or less,
has a structure in which carbides are dispersed so that the relationship C08/C10 S: 0.8 holds
between the carbide number Cos on the ferrite crystal grain boundary and the carbide number
Cro in the ferrite crystal grains, and has a cross-sectional hardness of 160 HV or less.
[0006]
Patent Document 3 discloses a medium- to high-carbon steel sheet with excellent
workability, wherein the sheet comprises C: 0.30 to 1.00 wt%, Si: 1.0 wt% or less, Mn: 0.2 to
- 2-
1.5 wt%, P: 0.02 wt% or less, and S: 0.02 wt% or less, has a structure in which carbides are
dispersed in ferrite so that the relationship CGs/CIG ~ 0.8 holds between the carbide number
CGB on the ferrite crystal grain boundary and the carbide number CIG in the ferrite crystal
grains, and simultaneously 90% or more of the total carbides are occupied by spheroidized
carbides having a long axis/short axis of2 or less.
[0007]
Patent Documents I to 3 describe that the greater the proportion of carbides in ferrite
grains, the more the workability is improved.
[0008]
In addition, Patent Document 4 discloses a steel sheet having excellent FB workability,
mold life, and cold formability after FB processing, wherein the sheet comprises C: 0.1 to 0.5
wt%, Si: 0.5 wt% or less, Mn: 0.2 to 1.5 wt%, P: 0.03 wt% or less, S: 0.02 wt% or less, has a
structure based on ferrite and carbide, and the amount Sgb of the carbide present on the ferrite
grain boundary is 40% or more, the above Sgb being defmed by Sgb = {Sm/(Son +Sin)} x 100
(wherein Son is the total area occupied by the carbides present on the grain boundary among
the carbides present per unit area and Sin is the total area occupied by the carbides present on
the grain boundary among the carbides present per unit area).
[0009)
However, in the technology described in Patent Document 1, annealing is performed at
a temperature of the Ac1 point or higher for softening in order to coarsen ferrite grain size and
carbide. But when annealing is performed at a temperature oftbe A c1 point or higher, rodlike/
plate-like carbides may precipitate during annealing. The carbides, even though capable of
reducing hardness, deteriorate workability, which is disadvantageous in terms of workability.
[0010]
The technologies described in Patent Documents 2 and 3 consider that the deterioration
of workability is caused by the low carbide spheroidization ratio of carbides precipitated on
the grain boundary, but do not take into account the problem of improving the spheroidization
ratio of grain boundary carbides. Techniques described in Patent Document 4 only specifY the
tissue factor, and Patent Document 4 does not discuss the relationship between workability
and mechanical properties.
[0011]
The technology described in Patent Document 5 is an invention made by focusing on
the relationship between tine blanking workability and the amount of carbide present in ferrite
grains and ferrite grain size. However, Patent Document 5 does not discuss what effect the
aggregate structure has on the plastic anisotropy.
[0012)
Patent Document 6 discloses a hot-rolled steel sheet in which the development of an
- 3-
aggregate structure otherwise developed by ro !ling is suppressed and a method for
manufacturing the sheet. However, Patent Document 6 does not discuss the relationship
between the aggregate structure otber than the aggregate structure developed by rolling and the
cold forgeability.
[0013]
The technology described in Patent Document 7 is an invention made by considering
that the hardness and the total elongation of a high-carbon hot-rolled steel sheet prior to
quenching are greatly influenced by the cementite density in the ferrite grains. The hot-rolled
steel sheet described in Patent Document 7 is characterized in that it has a microstructure
composed of ferrite and cementite, said microstructure having a cementite density ofO.lO
strips/J.!m2 or less in the ferrite grains. However, Patent Document 7 does not discuss what
effect the aggregate texture has on the plastic anisotropy.
[0014]
The technology described in Patent Document 8 is an invention made by considering
that the C,q value is related not only to mechanical properties and weldability but also to the
fatigue crack growth rate in steels having a fine structure. Patent Document 8 discloses that by
limiting the range of the Ceq value to a range of0.28% to 0.65%, the futigue resistance of the
steel material is improved and simultaneously weldability is secured. However, Patent
Document 8 does not discuss what effect the aggregate texture has on the plastic anisotropy.
Prior Art Documents
Patent documents
[0015]
[Patent Document 1] Japanese Patent No. 4465057
[Patent Document 2] Japanese Patent No. 4974285
[Patent Document 3] Japanese Patent No. 5197076
[Patent Document 4] Japanese Patent No. 5194454
[Patent Document 5] Japanese Unexamined Patent Publication No. 2007-270331
[Patent Document 6] Japanese Unexamined Patent Publication No. 2009-263718
[Patent Document 7] Japanese Unexamined Patent Publication No. 2015-17294
[Patent Document 8] Japanese Unexamined Patent Publication No. 2004-27355
Summary Of The Invention
Problem to be Solved by the Invention
[0016]
In view of the current state of the prior art, it is an object of the present invention to
address the problem of improving the cold workability of a steel sheet during forming, and to
provide a steel sheet that has solved the problem and a method for manufacturing the sheet.
Means to Solve the Problem
- 4-
[0017]
The present inventors have conducted intensive and extensive studies on methods for
solving the above-mentioned problems. As a result, the present inventors have found that by
controlling the dispersion state of the carbide in the structure of the steel sheet before cold
working through the optimization of the manufacturing conditions in the steps from hot rolling
to annealing, the carbide can be precipitated on the ferrite boundary and simultaneously the
aggregate structure in the hot rolled steel plate can be controlled, thereby leading to enhanced
cold workability.
[0018]
Further, we have found after intensive and extensive research that it is difficult to
manufacture a steel sheet that satisfies the above-mentioned conditions merely by devising hot
rolling conditions and annealing conditions separately, and that it can be manufactured by
optin1izing the above conditions in mutual cooperation in an integrated process of the hot
rolling and annealing steps.
[0019]
The present invention has been made based on the above fmdings, and the gist thereof
lies in:
[0020]
(I) A steel sheet having an excellent cold workability during forming, comprising, in
terms of% by mass:
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.0001 to 0.010%,
AI: 0.001 to 0.10%, and
a balance of Fe and inevitable impurities,
wherein (a) a ratio ofthe·number of carbides at a ferrite grain boundary relative to the
number of carbides in the ferrite grain is more than I,
wherein (b) a diameter of the ferrite grain is 5 1-1m or more and 50 flill or less,
wherein (c) an in-plane anisotropy ILI.rl of the r value standardized according to JIS Z 2254 is
0.2 or less,
wherein (d) a Vickers hardness of the steel sheet is 100 HV or more and 150 HVor less, and
wherein (e) a ratio of X-ray diffraction intensity ofthe {311) <011> orientation at the 112-
thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in the steel sheet is subjected to
X-ray diffraction is 3.0 or less.
- 5-
[0021]
(2) The steel sheet with excellent cold workability during forming described in the
above (1) further comprising, in terms of% by mass, one or a plurality of:
N: 0.0001 to 0.010%,
0: 0.0001 to 0.020%,
Cr: 0.00 I to 0.50%,
Mo: 0.001 to 0.10%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%,
Cu: 0.001 to 0.10%,
W: 0.001 to 0.10%,
Ta: 0.00 I to 0.1 0%,
Ni: 0.001 to 0.10%,
Sn: 0.001 to 0.050%,
Sb: 0.001 to 0.050%,
As: 0.001 to 0.050%,
Mg: 0.0001 to 0.050%,
Ca: 0.001 to 0.050%,
Y: 0.001 to 0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%, and
Ce: 0.001 to 0.050%.
[0022]
(3) A method for producing a steel sheet with excellent cold workability during
forming according to the above (I) or (2), the method comprising:
subjecting a steel strip having an ingredient composition according to claim I or 2 to
hot rolling by heating, followed by completing the finish hot rolling at a temperature range of
800°C or higher and 900°C or lower;
coiling the hot-rolled steel sheet at a temperature of 400°C or higher and 550°C or
lower;
pickling the hot-rolled steel sheet, and then subjecting the hot-rolled steel sheet to a
two-step type annealing in which the hot-rolled steel sheet is retained in two temperature
ranges~
wherein the two-step type annealing comprises
(i) subjecting the hot-rolled steel sheet to a frrst step annealing performed by retaining
said hot-rolled steel at a temperature range of650°C or higher and 720°C or lower for 3 hours
or longer and 60 hours or shorter, and then a second step annealing performed by retaining the
- 6-
hot-rolled steel at a temperature range of725°C or higher and 790°C or lower for 3 hours or
longer and 50 hours or shorter, and thereafter
(ii) cooling the hot-rolled steel sheet to 650°C or lower at a cooling rate of 1 °C/hour or
more and 30°C/hour or less.
[0023]
(4) The method for producing a steel sheet described in the above (3), wherein the steel
sheet has a cross-sectional shrinkage percentage of 40% or more.
Effect of the Invention
[0024]
According to the present invention, a steel sheet with excellent cold workability during
forming can be manufactured and provided.
Mode for Carrying Out the Invention
[0025]
A steel sheet with excellent cold workability during forming according to the present
invention (hereinafter may be referred to as "the inventive steel sheet") comprises, in terms of
%by mass:
C: 0.1 0 to 0.40%,
Si: 0.0 I to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.0001 to 0.010%,
AI: 0.001 to 0.10%, and
a balance of Fe and inevitable impurities,
the above sheet being characterized in that:
(a) the ratio of the number of carbides at a ferrite grain boundary relative to the number of
carbides in the ferrite grain exceeds I,
(b) the ferrite grain diameter is 5 J.lm or more and 50 J.lm or less,
(c) the in-plane anisotropy IMI of the r value standardized according to JIS Z 2254 is 0.2 or
less,
(d) the Vickers hardness is I 00 HV or more and 150 HV or less, and
(e) the ratio ofX-ray diffraction intensity of the {311} <0 11> orientation at the 1/2-thickness
portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample
with a random orientation distribution of crystal grains in the steel sheet is subjected to X-ray
diffraction is 3.0 or less.
[0026]
The method (hereinafter may be referred to as "the inventive method") ofthe present
invention for producing a steel sheet with excellent cold workability during forming is a
- 7-
method for producing the inventive steel sheet,
wherein a hot-rolled steel strip that has been obtained by subjecting a steel strip having
an ingredient composition ofthe inventive steel sheet to hot rolling by heating, followed by
completing the finish hot rolling at a temperature range of soooc or higher and 900°C or
lower, and by coiling the resulting hot-rolled steel sheet at a temperature of 400°C or higher
and 550°C or lower is, after pickling, subjected to two-step type annealing in which the sheet
is retained in two temperature ranges, whereupon
(i) the hot-rolled steel sheet is subjected to a first step annealing performed by retaining
said hot-rolled steel at a temperature range of 650°C or higher and noac or lower for 3 hours
or longer and 60 hours or shorter, and then subjected to a second step annealing performed by
retaining the hot-rolled steel at a temperature range of725°C or higher and 790°C or lower for
3 hours or longer and 50 hours or shorter, and thereafter
(ii) the sheet is cooled down to 650°C or lower at a cooling rate of I °C/hour or more
and 30°C/hour or less.
[0027]
Hereinafter, the inventive steel sheet and the inventive manufacturing method will be
described.
(0028]
First, the reasons for limiting the ingredient composition of the inventive steel sheet
will be described. The percentage relating to the ingredient composition means % by mass.
[0029]
C: 0.10 to 0.40%
C is an element that forms carbide in steel, and is effective for strengthening steel and
refrning ferrite grains. In order to prevent the surface of the steel sheet from being textured by
cold working and ensure the aesthetic appearance of surface of cold forged patis, it is
necessary to suppress the coarsening of ferrite grain size. However, when its content is less
than 0.1 0%, the volume fraction of the carbide is insufficient and the coarsening of carbides
during annealing cannot be suppressed. Therefore, Cis set to 0.10% or more, and preferably
0.12%ormore.
[0030]
On the other hand, when it exceeds 0.40%, the volume fraction of the carbide
increases, a large amount of cracks serving as fracture starting points are formed when a load
is instantaneously applied, and thus the impact resistance property decreases. Therefore, C is
set to 0.40% or less, and preferably 0.38% or less.
[00311
Si: 0.01 to 0.30%
Si is an element that acts as a deoxidizing agent and also affects the form of the
- 8-
carbide. In order to reduce the number of carbides in the ferrite grain and increase the number
of carbides on the ferrite grain boundaries, it is necessary to generate an austenite phase during
annealing in the two-step type annealing, and, after transiently dissolving the carbides, to cool
gradually to promote the precipitation of carbides at the ferrite grain boundaries.
[0032]
In the inventive steel sheet, the amount ofSi may preferably be as small as possible.
However, when it is reduced to less than 0.01 %, the manufacturing cost increases. Therefore,
Si is set to 0.01% or more.
[0033]
On the other hand, when it exceeds 0.30%, the ductility of ferrite lowers and breaking
may easily occur during cold working, resulting in reduced cold workability. Therefore, Si is
set to 0.30% or less, and preferably 0.28% or less.
[0034]
Mn: 0.30 to 1.00%
Mn is an element that controls the figuration of carbides in the two-step type annealing.
When its content is less than 0 .30%, it is difficult to precipitate carbides at the ferrite grain
boundaries in slow cooling after the second-step annealing. Therefore, Mn is set to 0.30% or
more, and preferably 0.33% or more.
[0035]
On the other hand, when it exceeds 1.00%, the hardness of ferrite increases and the
cold workability deteriorates. Therefore, Mn is set to 1.00% or less, and preferably 0.96% or
less.
[0036]
P: 0.0001 to 0.020%
P is an element that segregates at the ferrite grain boundaries and suppresses the
formation of grain boundary carbides. The amount ofP may preferably be as small as possible.
However, when Pis reduced to less than 0.0001% in the refming process, the refining cost
may greatly increase. Therefore, it is set to 0.0001% or more, and preferably 0.0013% or
more.
[0037]
On the other hand, when it exceeds 0.020%, the number percentage of the grain
boundary carbides decreases and the cold workability deteriorates. Therefore, P is set to
0.020% or less, and preferably 0.018% or less.
[0038]
S: 0.0001 to 0.010%
S is an element that fon11S a non-metallic inclusion such as MnS. Since a non-metallic
inclusion serves as the starting point for break generation during cold forging, the amount of S
- 9-
may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the
refming cost greatly increases. Therefore, S is set to 0.000 I% or more, and preferably
0.0012% or more.
[0039]
On the other hand, when it exceeds 0.010%, cold workability deteriorates. Therefore, S
is set to 0.010% or less, and preferably 0.007% or less.
[0040]
Al: 0.001 to 0.10%
Al is an element that acts as a deoxidizing agent for steel and stabilizes ferrite. When
its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Al is
set to 0.001% or more, and preferably 0.004% or more.
[0041 1
On the other hand, when it exceeds 0.1 0%, the number percentage of carbides on the
grain boundary decreases and the cold workability deteriorates. Therefore, Al is set to 0.10%
or less, and preferably 0.08% or less.
[00421
In addition to the above elements, the inventive steel sheet may contain one or a
plurality ofN: 0.0001 to 0.010%,0: 0.0001 to 0.020%, Cr: 0.001 to 0.50%, Mo: 0.001 to
0.10%, Nb: 0.001 to 0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta:
0.001 to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to
0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to
0.050%, La: 0.001 to 0.050%, and Ce: 0.001 to 0.050%, in order to improve the properties of
the inventive steel sheet.
[0043]
N: 0.0001 to 0.010%
N is an element that, when present in large amounts, causes the embrittlement of
ferrite. The amount ofN may preferably be as small as possible. However, when it is reduced
to less than 0.0001%, the refining cost greatly increases. Therefore, N should be 0.0001% or
more, and preferably 0.0006% or more. On the other hand, when it exceeds 0.010%, ferrite
embrittles and the cold forgeability deteriorates. Therefore, N should be 0.0 I 0% or less, and
preferably 0.007% or less.
[0044]
0: 0.0001 to 0.020%
0 is an element that, when present in large amounts, forms coarse oxides in steel. The
amount ofO may preferably be as small as possible. However, when it is reduced to less than
0.0001%, the refining cost increases greatly. Therefore, 0 is set to 0.0001% or more, and
preferably 0.0011% or more. On the other hand, when it exceeds 0.020%, coarse oxides are
- 10-
formed in the steel, the oxides serving as the starting point for break generation during cold
working. Therefore, 0 is set to 0.020% or less, and preferably 0.017% or less.
[0045]
Cr: 0.00 I to 0.50%
Cr is an element which enhances quenchability and contributes to the improvement of
strength and which is thickened to carbide and fonns stable carbide even in the austenitic
phase. When its content is less than 0.001%, the sufficient effect of improving quenchability
cannot be obtained. Therefore, Cr is set to 0.00 I% or more, and preferably 0.007% or more.
On the other hand, when it exceeds 0.50%, the carbide becomes stabilized thereby delaying
the dissolution of the carbide during quenching,and thus, it is feared that the desired quenching
strength may not be achieved. Therefore, Cr is set to 0.50% or less, and preferably 0.45% or
less.
[0046]
Mo: 0.001 to 0.10%
Like Mn, Mo is an element effective for controlling the figuration of carbides. When
its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Mo
is set to 0.001% or more, and preferably 0.010% or more. On the other hand, when it exceeds
0.1 0%, the in-plane anisotropy of the r value deteriorates and the cold workability deteriorates.
Therefore, Mo is set to 0.10% or less, and preferably 0.08% or less.
[0047]
Nb: 0.001 to 0.10%
Nb is an element which is effective for controlling the figuration of carbides and which
refmes the structure, thereby contributing to the enhancement of its toughness. When its
content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Nb
should be 0.00 I% or more, and preferably 0.004% or more. On the other hand, when it
exceeds 0.1 0%, a large number of fine Nb carbides precipitate, which leads to excessively
increased strength. It also causes the reduction in the number ratio of grain boundary carbides,
and the deterioration in cold forgeability. Therefore, Nb is set to 0.10 or less, and preferably
0.08% or less.
[0048]
V: 0.001 to 0.10%
Like Nb, V is an element which is effective for controlling the figuration of carbides
and which refines the structure, thereby contributing to the enhancement of its toughness.
When its content is less than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, Vis set to 0.001% or more, and preferably 0.004% or more. On the other hand,
when it exceeds 0.1 0%, a large number of fine V carbides precipitate, which leads to
excessively increased strength, to the reduced number ratio of grain boundary carbides, and to
- 11 -
the deteriorated cold forgeability. Therefore, V is set to 0.10 or less, and preferably 0.08% or
less.
[0049]
Cu: 0.001 to 0.10%
Cu is an element which segregates at the ferrite crystal grain boundary and forms fme
p~ecipitates thereby to contribute to the enhancement of strength. When its content is Jess than
0.00 I%, a sufficient effect of enhancing strength cannot be obtained. Therefore, Cu is set to
0.00 I% or more, and preferably 0.005% or more. On the other hand, when it exceeds 0. I 0%,
red heat embrittlement occurs and the productivity by hot rolling decreases. Therefore, Cu is
set to 0.10% or less, and preferably 0.08% or less.
[0050]
W: 0.001 to 0.10%
Like Nb and V, W is also an element effective for controlling the figuration of
carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, W is set to 0.001% or more, and preferably 0.003% or more. On the other hand,
when it exceeds 0.10%, a large number of fine W carbides precipitate, which leads to
excessively increased strength, to the reduced number ratio of grain boundary carbides, and to
the deteriorated cold forgeability. Therefore, W is set to 0.10 or less, and preferably 0.08% or
less.
[0051]
Ta: 0.001 to 0.10%
Like Nb, V and W, Ta is also an element effective for controlling the figuration of
carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, W is set to 0.00 I% or more, and preferably 0.005% or more. On the other hand,
when it exceeds 0.10%, a large number offme W carbides precipitate, which leads to
excessively increased strength, to the reduced number ratio of grain boundary carbides, and to
tbe deteriorated cold forgeability. Therefore, Ta is set to 0.10 or less, and preferably 0.08% or
less.
[0052]
Ni: 0.001 to 0.10%
Ni is an element effective for improving the toughness of patts. When its content is
less than 0.00 I%, a sufficient addition effect cannot be obtained. Therefore, Ni is set to
0.00 I% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.1 0%,
the number ratio of grain boundary carbides decreases and the cold forgeability deteriorates.
Therefore, Ni is set to 0.10% or less, and preferably 0.08% or less.
[0053]
Sn: 0.001 to 0.050%
- 12-
Sn is an element contaminated from a steel raw material (scrap). It segregates at the
grain boundary, leading to the decreased number ratio of grain boundary carbides. Therefore,
its content may preferably be as small as possible. However, when it is reduced to less than
0.001%, the refining cost will be greatly increased. Therefore, Sn is set to 0.001% or more,
and preferably 0.002% or more. On the other hand, when it exceeds 0.050%, ferrite embrittles
and cold forgeability deteriorates. Therefore, Sn is set to 0.050% or less, and preferably
0.040% or less.
[0054]
Sb: 0.001 to 0.050%
Like Sb, Sb is an element contaminated from a steel raw material (scrap). It segregates
at the grain boundary, leading to the decreased number ratio of grain boundary carbides.
Therefore, its content may preferably be as small as possible. However, when it is reduced to
less than 0.001%, the refming cost will be greatly increased. Therefore, Sb is set to 0.001% or
more, preferably 0.002% or more. On the other hand, when it exceeds 0.050%, the cold
forgeability deteriorates. Therefore, Sb is set to 0.050% or less, and preferably 0.040% or less.
[0055]
As: 0.001 to 0.050%
Like Sn and Sb, As is an element contaminated from a steel raw material (scrap). It
segregates at the grain boundary, thereby leading to a decrease in the number ratio of grain
boundary carbides. Therefore, its content may preferably be as small as possible. However,
when it is reduced to less than 0.001%, the refming cost increases greatly. Therefore, As is set
to 0.001% or more, and preferably 0.002% or more. On the other hand, when it exceeds
0.050%, the number ratio ofthe grain boundary carbides decreases and the cold forgeability
deteriorates. Therefore, As is set to 0.050% or less, and preferably 0.040% or less.
[0056]
Mg: 0.0001 to 0.050%
Mg is an element that can control the figuration of sulfides with the addition of its trace
amount. When its content is less than 0.0001%, a sufficient addition effect cannot be obtained.
Therefore, Mg is set to 0.000 I% or more, and preferably 0.0008% or more. On the other hand,
when it exceeds 0.050%, ferrite embrittles and the cold forgeability deteriorates. Therefore,
Mg is set to 0.050% or less, and preferably 0.040% or less.
[0057]
Ca: 0.001 to 0.050%
Like Mg, Ca is an element that can control the figuration of sulfides with the addition
of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be
obtained. Therefore, Ca is set to 0.001% or more, and preferably 0.003% or more. On the other
hand, when it exceeds 0.050%, coarse Ca oxides are formed, which serve as starting points of
- 13 -
break generation during cold forging. Therefore, Ca is set to 0.050% or less, and preferably
0.040% or less.
[0058]
Y: 0.001 to 0.050%
Like Mg and Ca, Y is an element that can control the figuration of sulfides with the
addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect
cannot be obtained. Therefore, Y is set to 0.001% or more, and preferably 0.003% or more. On
the other hand, when it exceeds 0.050%, coarse Y oxides are formed, which serve as starting
points of break generation dnring cold working. Therefore, Y is set to 0.050% or less, and
preferably 0.035% or less.
[0059]
Zr: 0.001 to 0.050%
Like Mg, Ca and Y, Zr is an element that can control the fignration of sulfides with the
addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect
cannot be obtained. Therefore, Zr is set to 0.001% or more, and preferably 0.004% or more.
On the other hand, when it exceeds 0.050%, coarse Zr oxides are formed, which serve as
starting points for break generation during cold working. Therefore, Zr is set to 0.050% or
less, and preferably 0.045% or less.
[0060)
La: 0.001 to 0.050%
La is an element that can control the figuration of sulfides with the addition of its trace
amount, but it is also an element that segregates at the grain boundary and causes a decrease in
tbe number ratio of grain boundary carbides. When its content is less than 0.001%, a sufficient
effect of controlling figuration cannot be obtained. Therefore, La is set to 0.001% or more, and
preferably 0.004% or more. On the other hand, when it exceeds 0.050%, the number ratio of
grain boundary carbides decreases and the cold workability deteriorates. Therefore, La is set to
0.050% or less, and preferably 0.045% or less.
[0061]
Ce: 0.001 to 0.050%
Like La, Ce is an element that can control the figuration of sulfides with the addition of
its trace amount, but it is also an element that segregates at the grain boundary and causes a .
decrease in the number ratio of grain boundary carbides. When its content is less than 0.00 I%,
a sufficient effect of controlling figuration cannot be obtained. Therefore, Ce is set to 0.001%
or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.050%, the
number ratio of grain boundary carbides decreases and the cold forgeability deteriorates.
Therefore, Ce is set to 0.050% or less, and preferably 0.045% or less.
[0062]
- 14-
The remainder of the ingredient composition of the inventive steel sheet is Fe and
unavoidable impurities.
[0063]
It is a novel fmding by the inventors that the inventive steel sheet has excellent cold
workability during forming, because, in addition to the above ingredient composition, it was
found, as a result of optimum hot rolling and annealing, that
(a) the ratio of the number of carbides at the ferrite grain boundary relative to the
number of carbides in the ferrite grain exceeds 1,
(b) the ferrite grain diameter is 5 J.lm or more and 50 J.lm or less,
(c) the in-plane anisotropy IMI of the r value standardized according to JIS Z 2254 is
0.2 or less,
(d) the Vickers hardness is 100 HV or more and 150 HV or less, and
(e) the ratio of X-ray diffraction intensity of the {311} <01 1 >orientation at the 112-
thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in the steel sheet is subjected to
X-ray diffraction is 3.0 or less.
[0064]
The above (a) to (e) will be described below.
[0065]
(a) The ratio of the number of carbides at the ferrite grain boundary relative to the number of
carbides in the ferrite grain exceeds 1:
The inventive steel sheet has a structure which is substantially composed of ferrite and
carbide, and in which the ratio of the number of carbides at the ferrite grain boundary relative
to the number of carbides in the ferrite grain exceeds I. Carbides are, in addition to cementite
(Fe3C) that is a compound of iron and carbon, compounds obtained by replacing Fe in
cementite with an element such as Mn and Cr, and alloy carbides (M23C6, M6Co, MC, etc.,
wherein M is Fe and another additive metal element).
[0066]
When a steel sheet is formed into a predetermined part shape, a shear band is formed in
the macrostructure of the steel sheet, and slip deformation is generated and concentrated in the
vicinity of the shear band. The slip deformation involves propagation of dislocations, and
regions with high dislocation density are formed in the vicinity of the shear band. As the strain
amount applied to the steel sheet increases, the slip deformation is promoted and thereby the
dislocation density increases. In cold forging, strong processing exceeding an equivalent strain
of 1 is applied.
[0067]
For this reason, in the conventional steel sheet, generation of voids and/or cracks due
- 15 -
to the increased dislocation density could not be prevented, and it was difficult to improve cold
forgeability.
[0068]
In order to solve the above challenging problems, it is effective to suppress the
formation of shear bands during forming. From the viewpoint of a microstructure, shear band
formation is a phenomenon in which a slip generated in one crystal grain crosses the crystal
grain boundary and propagates continuously to an adjacent crystal grain. Therefore, in order to
suppress the formation of a shear band, it is necessary to prevent the propagation of slippage
beyond crystal the grain boundary.
(0069]
Carbides in the steel sheet are tenacious particles that hinder slippage. Therefore, the
presence of carbides at the ferrite grain boundaries would make it possible, for the frrst time, to
suppress the formation of a shear band and thereby to improve cold forgeability.
[0070]
Based on the theory and principle, it is considered that cold forgeability is strongly
influenced by the coverage rate of carbides at the ferrite grain boundaries. Therefore, it
becomes necessary to measure the coverage rate with high accuracy.
[0071]
In order to measure the coverage rate of carbides at the ferrite grain boundaries in a
three-dimensional space, serial sectioning SEM observation or repeated three-dimensional
EBSP observation become essential in which sample cutting by FIB and observation are
repeated in the scanning electron microscope. However, these methods take a huge amount of
measurement time and the accumulation of technical know-how becomes indispensable. We
clarified this fact and concluded that common analytical methods are not suitable.
(0072]
Therefore, as a result of searching a simple and highly accurate evaluation index, the
present inventors have found that cold forgeability can be evaluated by using, as an index, the
ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides
in the ferrite grain, and that cold forgeability can be remarkably improved when the ratio of
the number of carbides at the ferrite grain boundary relative to the number of carbides in the
ferrite grain is more than 1.
(0073]
Any of buckling, folding and convolution of a steel sheet that occurs during cold
working is caused by the localization of strain accompanying the formation of a shear band.
Therefore, by allowing the carbide to exist at the ferrite grain boundaries, the formation ofthe
shear band and the localization of strain can be alleviated and the generation of buckling,
folding and convolution can be suppressed.
- 16 -
[0074]
When the spheroidization percentage of carbides on the crystal grain boundary is less
than 80%, strains are concentrated locally on the rod-shaped or plate-shaped carbides, and
voids and/or cracks are likely to occur. Therefore, the carbide spheroidization ratio on the
crystal grain boundary may preferably be 80% or more, and more preferably 90% or more.
[0075]
When the average particle diameter of the carbide in the ferrite grain and the carbide at
the ferrite grain boundaries is less than 0.1 J.Lill, the hardness of the steel sheet remarkably
increases and the workability deteriorates. Therefore, the average particle diameter of the
carbide may preferably be 0.1 J.Lm or more, and more preferably 0.17 J.Lm or more. On the
other hand, when the average particle diameter of the carbide exceeds 2.0 J.Lm, fissures occur
with the coarse carbide serving as a starting point during cold working, and thus the cold
workability deteriorates. Therefore, the average particle diameter of the carbide may
preferably be 2.0 J.Lm or less, and more preferably 1.95 J.Lm or less.
[00761
Subsequently, the method of observing and measuring the structure will be described.
[0077]
Observation of the carbide is carried out by a scanning electron microscope. Prior to
observation, samples for structure observation are polished by wet polishing with emery paper
and polishing with diamond abrasive grains having an average particle size of I J.Lm. After
polishing the observation surface to a mirror finish, the structure is etched with a 3% nitric
acid-alcohol solution.
[0078]
Among the magnification for observation, within 3000 times, a magnification capable
of discriminating between ferrite and carbide is selected. At the selected magnification, eight
images with a viewing field of 30 J.Lm x 40 J.Lm are randomly photographed at the 1/4 plate
layer thickness.
[0079]
With respect to the tissue image obtained, the area of each carbide contained in the
region is measured in detail by an image analysis software represented by Mitsuya Shoji Co.
Ltd. (Win ROOF). A circle equivalent diameter(= 2X'i(area/3.14)) is obtained from the area of
each carbide, and the average value is taken as the carbide particle diameter.
[0080]
Fmiber, the spheroidization ratio of the carbide was determined by approximating the
carbide to an ellipse having an equal area and equal moment of inertia, and then by calculating
the proportion of the carbides in which the ratio of the maximum length to the maximum
length in the perpendicular direction is less than 3.
- 17-
[0081]
In order to suppress the effect of measurement error due to noise, carbides having an
area ofO.Ol J.Lm2 or more among the carbides in grains and grain boundaries were counted and
the carbides having an area of 0.0 I J.Lm2 or less were excluded from evaluation.
[0082]
The number of carbides present on the ferrite grain boundary was counted, and from
the total number of carbides the number of carbides in the ferrite grain was determined by
subtracting the number of carbides on the ferrite grain boundary. Based on the measured
number, the ratio of the uumber of carbides on the grain boundary relative to the number of
carbides in the ferrite grain was determined.
[0083]
(b) The ferrite grain diameter is 5 J.Lm or more and 50 J.Lm or less:
In the structure after annealing the cold rolled steel sheet, the cold workability can be
improved by setting the ferrite grain diameter to 5 J.Lm or more. When the ferrite grain size is
less than 5 J.lill, the hardness increases and fissures and cracks tend to generate easily during
cold working. Therefore, the ferrite grain size is set to 5 J.Lm or more, and preferably 7 J.Lm or
more.
[0084]
On the other hand, when it exceeds 50 J.!ill, the number of carbides on the crystal grain
boundary that suppress slippage propagation decreases and the cold workability deteriorates.
Therefore, that the ferrite grain size is set to 50 J.Lm or less, and preferably 37 J.Lm or less.
[0085]
TI1e ferrite grain diameter is measured in the above-described polishing method,
wherein the observation surface of the sample is polished to a mirror surface, followed by
etching with a 3% nitric acid-alcohol solution. The structure of the observation surface is then
examined with an optical microscope or a scanning electron microscope, and a line segment
method is then applied to the image photographed to determine the ferrite grain diameter.
[0086]
(c) The in-plane anisotropy ILl.rl ofthe r value standardized according to JIS Z 2254 is
0.2 or less:
The in-plane anisotropy ILl.rl of the plastic strain ratio (r value) of the steel sheet is
measured in a method in accordance with ITS Z 2254. The r value (0° direction: r0, 45°
direction: r45, 90° direction: r90) measured by taking test strips from each direction of0°
direction, 45° direction and 90° direction with respect to the rolling direction was used to
calculate the following equation.
ILl.rl ~ (ro - 2r4s + r9o)/2
[0087]
- 18-
By setting the in-plane anisotropy IL'lrl of the plastic strain ratio (r value) of the steel
sheet to 0.2 or less, the cold workability can be improved. When IL'lrl exceeds 0.2, the thickness
of parts and the height of the earing become nneven during drawing. Therefore, the in-plane
anisotropy IMI is set to 0.2 or less.
[0088]
(d) The Vickers hardness is 100 HV or more and !50 HV or Jess:
By setting the Vickers hardness of the steel sheet to 100 HV or more and !50 HV or
less, the cold workability can be improved. When the Vickers hardness is less than I 00 HV,
buckling can easily occur during cold working. Therefore, the Vickers hardness is set to 100
HV or more, and preferably II 0 HV or more.
[0089]
On the other hand, when the Vickers hardness exceeds !50 HV, the ductility decreases
and the internal breaking tends to occur easily during cold forging. Therefore, the Vickers
hardness is set to !50 HV or Jess, and preferably 146 HV or less.
[0090]
(e) The ratio of X-ray diffraction intensity of the {31 I} orientation at the 1/2-
thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in the steel sheet is subjected to
X-raydiffraction is 3.0 or less:
In cold forging, in addition to controlling the figuration of carbides, the draw
formability during cold forging must be secured. In order to improve the draw formability
during cold forging, plastic anisotropy such as in-plane anisotropy IMI must be improved. For
that purpose, the aggregate structure of a hot-rolled steel sheet must be controlled. For
evaluation of the aggregate structure, analysis by X-ray diffraction on a plane parallel to the
plate surface at the 1/2 thickness portion of the hot-rolled steel plate is used.
[0091]
One surface of a hot-rolled steel plate is ground to a 112 plate thiclmess surface in
parallel to the surface to expose a 1/2 plate thickness surface, followed by the analysis of the
1/2 plate thickness surface by X-ray diffraction. As the X-ray diffraction, X-ray diffraction by
Mo bulb may be used. Diffraction intensities of diffraction orientations {110}, {220}, {211}
and {31 0} by reflection are obtained, and based thereon, the orientation distribution function
(ODF) is created.
[0092]
The X-ray diffraction intensity ratio is determined by using the diffraction intensity
data of the 1/2 plate thickness surface obtained from the ODF and the diffi·action intensity data
of random orientation of the hot-rolled steel sheet. Specifically, as a standard sample in which
the metallic structure has no accumulation in a specific direction, a sample obtained by
- 19-
sintering powder iron of a hot-rolled steel sheet to be measured or the powder before sintering
is used to determine the diffraction intensity under the same conditions as when the diffraction
intensity data of the 112 plate thickness surface was obtained. The part to be collected as the
standard sample is not particularly limited and may be any part of the hot-rolled steel sheet.
The X-ray diffraction intensity ratio in a specific orientation is a numerical value obtained by
dividing the diffraction intensity in the specific direction of the 112 plate thickness surface
obtained from the ODF by the diffraction intensity of the standard sample.
[0093]
When the X-ray diffraction intensity ratio of the {311} <011 >orientation obtained by
the above-described ODF analysis is set to II, it is necessary that this II is 3.0 or less, and
preferably 2.5 or less for the random aggregate structure during hot rolling. When a random
aggregate structure having I1 of 3.0 or less can be obtained, the plasticity anisotropy is
reduced and the cold formability is improved.
[0094]
Next, the inventive manufacturing method will be described.
[0095]
The manufacturing method according to the present invention is characterized in that
the hot rolling and the annealing are consistently managed to control the structure. After
continuously casting a steel strip having a predetennined ingredient composition, the steel
strip is subjected to hot rolling by heating to complete fmish hot rolling at a temperature range
of 800°C or higher to 900°C or lower, coiled at 400°C or higher and 550°C or lower to obtain
a hot-rolled steel sheet. The hot-rolled steel sheet is, after pickling, subjected to a two-step
type annealing in which the hot-rolled steel sheet is maintained in two temperature ranges,
whereupon
(i) the hot-rolled steel sheet is subjected to a first step annealing performed by retaining
said hot-rolled steel at a temperature range of 650°C or higher and 720°C or lower for 3 hours
or longer and 60 hours or shorter, and then subjected to a second step annealing performed by
retaining the hot-rolled steel at a temperature range of725°C or higher and 790°C or lower for
3 hours or longer and 50 hours or shorter, and thereafter
(ii) the hot-rolled steel sheet is cooled to 650°C or lower at a cooling rate of I °C/hour
or more and 30°C/hour or less,
and thus a steel sheet excellent in cold workability during fom1ing can be produced.
[0096]
By the hot rolling and annealing mentioned above, a structure composed of fine
pearlite and bainite can be formed as the structure of the steel sheet.
[0097]
The processing conditions will be described below.
-20-
[0098]
Heating temperature of a steel strip: 1000°C or higher and 1250°C or lower
The heating temperature of the steel strip subjected to hot rolling may preferably be
1 ooooc or higher and 1250°C or lower, and the heating time may preferably be 0.5 hour or
longer and 3 hours or shorter.
[0099]
When the heating temperature is lower than 1 000°C or the heating time is shorter than
0.5 hour, the microsegregation and/or macro segregation formed by casting are not eliminated,
and regions in which Si, Mn, etc., are locally concentrated inside the steel material may
remain, and thus the impact resistance property of the steel material is lowered. Therefore, the
heating temperature may preferably be 1 ooooc or higher, and preferably 0.5 hour or longer.
[0100]
On the other hand, when the heating temperature exceeds 1250°C or the heating time
exceeds 3 hours, decarburization from the surface layer of the steel strip becomes conspicuous,
and austenite grains in the surface layer grow abnormally during heating before carburizing
and quenching, and the impact resistance property of the steel strip is deteriorated. Thus the
heating temperature may preferably be 1250°C or lower, and the heating time may preferably
be 3 hours or shorter.
[0101]
Finish hot rolling temperature: 800°C or higher and 900°C or lower
Finish hot rolling is completed at soooc or higher and 900°C or lower. When the finish
hot rolling temperature is lower than 800°C, the deformation resistance of the steel strip
increases, the rolling load increases markedly, the wear amount of the roll increases, and the
productivity decreases. Therefore, the fmish hot rolling temperature is set to 800°C or higher,
and preferably 820°C or higher.
[0102]
On the other hand, when the finish hot rolling temperature exceeds 900°C, thick scales
are generated during plate passing on the ROT (Run Out Table), scratches are generated on the
surface of the steel sheet due to the scale, and cracks are generated starting from scratches
when an impact load is applied after cold forging and carburizing and annealing, leading to
reduced impact resistance property of the steel sheet. Therefore, the finish hot rolling
temperature is set to 900°C or lower, and preferably 880°C or lower.
[0103]
Cooling rate on ROT: 1 ooC/sec or more and l 00°C/sec or less
The cooling rate at the time of cooling the hot-rolled steel sheet on the ROT after finish
hot rolling may preferably be 1 0°C/sec or more and 1 00°C/sec or less. When the cooling rate
- 21 -
is less than 10°C/sec, thick scales are generated during cooling and the occurrence of scratches
on the surface of the steel sheet due to the scales cannot be suppressed. Therefore, the cooling
rate is set to I 0°C/sec or more, and more preferably 20°C/sec or more.
[0104]
On the other hand, when the cooling rate exceeds I 00°C/sec, the steel sheet is cooled at
a cooling rate exceeding 100°C/sec from the surface layer to the inside of the steel sheet, the
outermost layer part ofthe steel sheet is excessively cooled, and a low-temperature
transformed structure such as bainite or martensite is formed.
(0105]
At the time of discharging the hot-rolled coil cooled from 100°C to room temperature
after coiling, microcracks are generated in the low-temperature transformed structure. It is
difficult to remove the microcracks in the subsequent pickling step and cold rolling step, and
fissures progress from the microcracks as a starting point during cold working, leading to
reduced cold workability. Therefore, the cooling rate may preferably be 1 00°C/sec or less.
[0106]
Note that the above cooling rate refers to the cooling capacity from the cooling facility
at each water injection zone from the point at which the hot-rolled steel sheet after the fmish
hot rolling is cooled at the water injection zone after passing through the water-free zone to a
point at which it is cooled to the coiling target temperature on the ROT, and does not refer to
the average cooling rate fiom the water injection starting point to the temperature at which it is
coiled by the coiling device.
(0107]
Coiling temperature: 400°C or higher and 550°C or lower
The coiling temperature is set to 400°C or higher and 550°C or lower. When the
coiling temperature is lower than 400°C, the austenite which was not transformed before
coiling is transformed into hard martensite, cracks are generated in the surface layer of the
steel sheet during discharge of the hot-rolled coil, leading to reduced workability. Therefore,
the coiling temperature is set to 400°C or higher, and preferably 430°C or higher.
[0108]
On the other hand, when the coiling temperature exceeds 550°C, pearlite having a large
lamellar spacing is generated and thick needle-shaped carbides having high thermal stability
are formed, and even after the two-step type annealing, needle-shaped carbides remain. Since
fissures are generated during cold working with these needle-shaped carbides as a starting
point, the coiling temperature is set to 550°C or lower, and preferably 520°C or lower.
[0109]
The hot-rolled coil manufactured under the above conditions is annealed, after
pickling, in a two-step type annealing which retains the coil in two temperature ranges. The
- 22-
first-step annealing and the second-step annealing may be either box annealing or continuous
annealing. By controlling the stability of carbides by the two-step type annealing, the
formation of carbides on the ferrite grain boundary and the spheroidization ratio of carbides on
the ferrite grain boundary can be enhanced.
[0110]
The two-step type annealing will be described below.
[0111]
The first step annealing is carried out in a temperature range of the Ac1 point or lower
to coarsen carbides and enrich alloy elements to increase the thermal stability of carbides.
Thereafter, the temperature is raised to a range from Ac1 point or higher to A3 point or lower
to generate austenite in the structure.
(0112]
Thereafter, by gradual cooling, the austenite is transformed into ferrite and the carbon
concentration in the austenite is increased. By proceeding slow cooling, carbon atoms are
adsorbed to the carbides remaining in the austenite, and thus the carbide and austenite come to
cover the grain boundary of the ferrite. Finally it becomes possible to form a structure in
which many spheroidized carbides are present in the grain boundary of the ferrite.
[0113]
When the residual carbides are small in quantity while maintaining the temperature
range of Ac1 point or higher to A3 point or lower, pearlite, rod-shaped carbides and plate-like
carbides are produced during cooling. When these pearlite, rod-shaped carbides and plate-like
carbides are produced, the workability of the steel sheet is remarkably deteriorated. Therefore,
increasing the number of residual carbides in the temperature range from Ac1 point or higher
to A3 point or lower is an important factor to enhance the workability of the steel sheet.
[0114]
By using a steel sheet structure obtained under the above hot rolling condition, the
thermal stability of carbides at a temperature of Ac1 point or lower can be secured. Therefore,
an increase in the number of residual carbides in the temperature range from Ac1 point or
higher to A3 point or lower can be targeted.
[0115]
Hereinafter, an annealing condition for the two-step type annealing will be described.
[0116]
First step annealing
Temperature range: 650°C or higher and 720°C or lower
Retention time: 3 hours or longer and 60 hours or shorter
In the first step annealing, the annealing temperature is set to 650°C or higher and
nooc or lower. When the annealing temperature of the first step is lower than 650°C, the
- 23 -
stability of the carbide becomes insufficient and it becomes difficult to allow the carbide to
remain in the austenite in the second step annealing. Therefore, the temperature of the first
step annealing is set to 650°C or higher, and preferably 670°C or higher.
[0117]
On the other hand, when the temperature of the first step annealing exceeds 720°C,
austenite is generated before enhancing the stability of the carbide, which makes it difficult to
control the required change in the structure. Therefore, the first step annealing temperature is
set to 720°C or lower, and preferably 700°C or lower.
(0118]
The retention time at the first step is 3 hours or longer and 60 hours or shorter. When
the retention time is lower than 3 hours, the stability of the carbide is insufficient and it
becomes difficult to allow the carbide to remain at the second step annealing. Therefore, the
retention time of the first step is set to 3 hours or longer. On the other hand, when the retention
time of the first step exceeds 60 hours, improvement of the stability of the carbide cannot be
expected and furthermore the productivity is lowered. Therefore, the retention time of the first
step is set to 60 hours or shorter, and preferably 55 hours or shorter.
[0119]
The annealing atmosphere is not limited to a specific atmosphere. For example, it may
be either a nitrogen atmosphere having a nitrogen content of95% or more, a hydrogen
atmosphere having a hydrogen content of95% or more, or an atmospheric atmosphere.
[0120]
Second step annealing
Temperature range: 725°C or higher and 790°C or lower
Retention time: 3 hours or longer and 50 hours or shorter
In the second step annealing, the annealing temperature is set to nsoc or higher and
790°C or lower. When the second-step annealing temperature is lower than 725°C, the amount
of austenite produced is small and the number ratio of carbides on the ferrite grain boundary is
lowered. Therefore, the second-step annealing temperature is set to 725°C or higher, and
preferably 745°C or higher.
[0121]
On the other hand, when the second-step annealing temperature exceeds 790°C, it
becomes difficult to allow the carbide to remain in the austenite and to control the required
structure change. Therefore, the second-step annealing temperature is set to 790°C or lower,
and preferably 770°C or lower.
[0122]
The retention time of the second step is set to 3 hours or longer and 50 hours or shorter.
When the retention time ofthe second step is less than 3 hours, the amount of austenite
- 24-
produced is small, dissolution of the carbide in the ferrite grains is insufficient, and it becomes
difficult to increase the number ratio of carbides on the ferrite grain boundary. Therefore, the
retention time ofthe second step is set to 3 hours or longer, and preferably 5 hours or longer.
[0123]
On the other hand, when the retention time ofthe second step exceeds 50 hours, it
becomes difficult to allow the carbide to remain in the austenite. Therefore, the retention time
of the second step is set to 50 hours or shorter, and preferably is 46 hours or shorter.
[0124]
The annealing atmosphere is not limited to a specific atmosphere. For example, it may
be either a nitrogen atmosphere having a nitrogen content of95% or more, a hydrogen
atmosphere having a hydrogen content of95% or more, or an atmospheric atmosphere.
[0125]
After completion of the two-step type annealing, the hot-rolled steel sheet is cooled,
whereupon it is cooled to 650°C at a cooling rate of 1 °C/hour or more to 30°C/hour or less.
[0126]
Cooling rate to a temperature of 650°C or lower: 1 °C/hour or more and 30°C/hour or
less
Since the temperature range for controlling the structure change by slow cooling is
sufficient up to 650°C, it is only necessary to control the cooling rate in the temperature range
up to 650°C. After reaching a temperature of 650°C or lower, it may be cooled to room
temperature within the above range without controlling the cooling rate.
[0127]
It may be preferable that the cooling rate is slow in order to gradually cool the austenite
produced in the second step annealing to transform into ferrite and allow carbon to be
adsorbed to the carbides remaining in the austenite. However, when the cooling rate is less
than I °C/hour, the time required for cooling increases and the productivity decreases.
Therefore, the cooling rate is 1 °C/hour or more, and preferably 5°C/hour.
[OWl]
On the other hand, when the cooling rate exceeds 30°C/hour, austenite transforms to
pearlite, the hardness of the steel sheet increases, the cold forgeability deteriorates, and the
impact resistance property of the steel sheet after carburizing quenching and tempering
decreases. Therefore, the cooling rate is set to 30°C/hour or less, and preferably 26°C/hour or
less.
[0129]
Further, according to the inventive production method, a steel sheet with excellent cold
workability during forming can be produced in which the ingredient composition is, in terms
of% by mass, comprising: C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00%, P: 0.0001
- 25 -
to 0.020%, S: 0.0001 to 0.010%, and AI: 0.001 to 0.1 0%, the balance being Fe and
unavoidable impurities, the metal structure is substantially composed of ferrite and
spheroidized carbides, and (a) the ratio of the number of carbides at the fen-ite grain boundary
to the number of carbides in the fen-ite grain exceeds I, (b) the ferrite grain size is 5 1-1m or
more and 50 1-1m or less, (c) the in-plane anisotropy IMI of the r value standardized according
to JIS Z 2254 is 0.2 or less, (d) the Vickers hardness is 100 HVor more and 150 HV or less,
the cross-sectional shrinkage percentage is 40% or more, and the ratio of X-ray diffraction
intensity of the {311} <0 II> orientation at the 1/2-thickness portion of the steel sheet relative
to the X-ray diffraction intensity obtained when a sample having the random orientation
distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less.
[0130]
The cross-sectional shrinkage percentage is defined by tbe following formula (1 ). A
large value of this ratio means that the local defonnability is high, and as the value of the
formula (I) increases, the workability of the steel sheet increases.
Sectional shrinkage percentage (%) = I 00 - (cross-sectional area at tensile fracture/initial
cross-sectional area) x 100 - - -Equation (I)
[01311
As described above, the present invention is characterized in that by rolling control and
heat treatment after rolling, a structure in which carbides (that is, cementite) are uniformly
dispersed is formed, so that the crystal anisotropy can be eliminated. Therefore, in the present
invention, the random intensity ratio of the {311} <0 11 > orientation at the 112 plate thickness
portion of the steel sheet can be made 3.0 or less.
[Examples]
[0132]
Next, examples will be described, but the level of examples is an example of
conditions adopted for confirming the feasibility and effectiveness of the present invention,
and the present invention is not limited to this one condition example. The present invention
can adopt various conditions as long as the object ofthe present invention is achieved without
departing from the gist of the present invention.
[0133]
[Example I]
In order to investigate the effect of hot rolling conditions, a continuous cast strip (steel
ingot) having the ingredient composition shown in Table 1 was subjected to hot rolling under
the conditions shown in Table 2 to produce a hot-rolled coil having a thickness of3.0 mm.
Incidentally, the steel type described as "Developed steel" in the column of"Remarks" in
Table 1 has a composition included in the composition range of the steel sheet according to the
present invention. Also, the steel type described as "Comparative steel" in the column of
-26-
"Remarks" in Table I has a composition outside the composition range of the steel sheet
according to the present invention. In addition, the ingredients that do not satisfy the
composition conditions of the steel sheet according to the present invention are underlined.
[0134]
A sample for characterization was prepared as follows: a hot-rolled coil, after pickling,
was placed in a box-type annealing furnace, the atmosphere was controlled to 95% hydrogenS%
nitrogen, the coil was heated from room temperature to 705°C and was retained for 36
hours to make the temperature distribution uniform in the hot-rolled coil. The coil was then
heated to 760°C and retained at 760°C for I 0 hours, and was then cooled to 650°C at a cooling
rate of I 0°C/hour, then furnace-cooled to room temperature to prepare the sample for
characterization. The structure of the sample was measured by the method described above.
[0135]
[Table I]
c Si Mn p s A1 N 0 Cr Mo
A 0.21 0.13 0.53 0,0048 0.0084 0.618 0.0053 0.098
B 0.24 0.28 0.91 0.0039 0.0018
c 0.34 0.19 0.51 0.0133 0.0015
D 0.18 0.18 0.50 0.0151 0.0092 0,0081 0.461
E 0.32 0.03 0.73 0.0149 0.0094 0.018 0.0018 0.087
F 0.22 0.06 0.98 0.0144 0.0088 0.082
G 0.17 0.13 0,94 0.0199 0.0026 0.038 0.277
H 0.37 0.11 0.38 0.0144 0.0087
! 0.33 0.10 0.96 0.0010 0.0019 1.180 0.480
J 0.29 0.16 0.81 0.0132 0.0031
K 0.11 0.02 0.41 0.0117 0.0360
L 0.25 0.11 0,63 0.0167 0.0058 0.0007 0.090
M 0.19 1.21 0.91 0.0027 0.0084
N 0.31 0.23 1.27 0.0010 0.0019
0 0.51 0.14 0,65 0.0059 0.0013
p 0.20 0.13 0.92 0.0153 0.0091
Q 0.11 0.07 0.64 0.0060 0.0051 0.0013 0.0131
R 0.27 0.19 0.70 0.0185 0.0033
s 0.27 0.13 0.80 0.0044 0.0002 0.071
T 0.25 0.14 0.38 0.0142 0.0097
u 0.12 0.25 0,86 0.0154 0.0045
w 0.40 0.08 0.78 0.0184 0.0087 0.026 0.025
X 0.13 0.20 0.96 0_0061 0.0034
y 0.21 0.21 0.41 0.0139 0.0038 0.092
z 0.11 0.05 0.59 0.0195 0.0010
AA 0.13 0.09 0,76 0.0127 0.0079
AB 0.19 0.22 0.89 0.0005 0.0021
AC 0.12 0.02 0.63 0.0168 0.0098
AD 0.30 0.28 0.71 0.0169 0.0031 -
Nb v Cu w Te Ni Sn Sb A'
0.084 0.004 0.015
0.049 0.083 0.018
0.017
0.080 0.020 0.034
0.073 0.039
0.037
0.020 0.012
0_037
0,033 0035
0.012 0.080
0.060 0.016
0.088 0.055
0.093
0.018 0.035 L_ ... ---"~·--- ---- 0.041
------
Mg Ca y Zr La
0.0149
0.0459
0.018
0.021
0.014
0.046 0.048
0.0423 0.035
0.046 0.048
0.037
0.0280 0.036
0.038
0.039
c,
0.045
0.008
0.015
0.030
Remarks
Com arative steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Comparative steel
Developed steel
Comparative steel
Developed steel
ComJ?arative steel
ComQarative steel
ComQarative steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
Developed steel
N..._ ,
[Table 2]
Hot rolling Condition
Ferrite
Finish hot
Coiling
Carbide
grain Vickers
rolling diameter hardness
temp, diameter
temp. [OC] ~m] [!illl]
[HV]
[OCJ
A-1 891 505 0.91 18.5 125
B-1 832 500 1.17 20.7 124
C-1 850 489 1.08 20.7 120
D-1 742 532 0.99 21.5 111
E-1 811 617 0.95 16.8 116
F-1 937 512 1.18 21.1 112
G-1 871 409 1.27 23.9 110
H-1 837 522 0.99 20.7 117
1-1 802 484 1.03 16.9 121
J-1 841 407 1.28 22.5 116
K-1 871 475 1.11 29.7 104
L-1 898 443 1.20 22.7 110
M-1 873 467 1.11 18.1 172
N-1 850 414 1.37 22.8 152
0-1 834 413 1.22 20.0 127
P-1 856 485 1.22 22.8 111
Q-1 873 428 1.25 29.3 116
R-1 845 348 1.32 24.0 114
S-1 837 454 1.23 21.9 114
T-1 830 671 0.75 17.4 116
U-1 881 419 1.29 27.6 112
W-1 853 535 1.09 18.5 121
X-1 885 477 1.24 25.8 112
Y-1 874 436 1.14 25.7 110
Z-1 899 485 1.16 27.7 108
AA-1 895 423 1.28 27.6 102
AB-1 855 434 1.27 23.9 115
AC-1 895 381 1.30 30.1 109
AD-1 854 463 1.17 20.7 125
Grain
Cross-sectional
{311)<011>
boundary
shrinkage
X-ray
carbide diffraction
percent
No. I grain [%] intensity
carbide No. ratio (Il)
3.7 51.0 5.5
6.7 49.7 1.6
4.6 51.3 2.1
6.3 54.0 5.6
5.4 29.7 1.5
7.2 52.8 2.5
8.7 56.3 1.5
3.7 55.7 2.0
7.1 49.5 5.0
6.9 43.9 1.9
4.4 31.4 1.4
6.9 54.7 2.3
5.4 34.9 4.6
9.3 38.5 2.4
5.7 38.4 2.0
7.2 57.1 2.2
6.2 54.1 2.0
6.5 51.1 2.2
7.1 56.6 1.8
2.8 38.7 1.4
7.3 59.0 2.7.
6.1 55.! 1.4
7.5 58.0 2.9
5.8 44.8 2.0
5.5 57.5 1.6
6.9 42.6 2.4
7.3 53.3 2.5
6.4 54.4 1.9
5.9 52.3 1.4
In-plane
anisotropy
lllrl
0.44
0.05
0.10
0.48
0.06
0.18
0.04
0.10
0.39
0.09
0.06
0.12
0.37
0.14
0.12
0.11
0.12
0.14
0.08
0.05
0.19
0.05
0.!9
0.12
0.06
0.14
0.16
0.11
0.06
Remarks
Com:Qarative steel
Inventive steel
Inventive steel
Comnarative steel
Comnarative steel
Comnarative steel
Inventive steel
Inventive steel
Comnarative steel
Inventive steel
Com12atative steel
Inventive steel
Comnarative steel
Comnarative steel
Comnarative steel
Inventive steel
Inventive steel
Com12arative steel
Inventive steel
Com)2arative steel
Inventive steel
Inventive steel
Inventive steel
Inventive steel
Inventive steel
Inventive steel
Inventive steel
Comgarative steel
Inventive steel
N
00
- 29-
[0136]
The cold workability was evaluated using the notched tensile test and the in-plane
anisotropy of the r value. In the notched tensile test, a notched tensile test strip was taken from
an as-annealed material with a thickness of3 mm, and a tensile test was performed in the
rolling direction to determine the cross-sectional shrinkage percentage, and the local
deformability was evaluated. When the cross-sectional shrinkage percentage is 40% or more, it
was rated as superior.
[0137]
Further, the in-plane anisotropy of the r value was rated as superior when the in-plane
anisotropy ILI.rl of the r value standardized according to JIS Z 2254 of an as-annealed material
with a thickness of3 mm was 0.2 or less.
[0138]
In order to determine the X-ray diffraction intensity ratio (II) of {311} <011>, X-ray
diffraction with an Mo tube was performed from the center of the plate thickness of each
sample followed by an ODF analysis. Based on the results obtained by the ODF analysis, the
II was determined.
[0139]
Table 2 shows, for each of the samples prepared, the results of the carbide diameter,
the ferrite grain diameter, the Vickers hardness, the ratio of the number of carbides at the
ferrite grain boundary relative to the number of carbides in the ferrite grain, the cross-sectional
shrinkage percentage, the X-ray diffraction intensity ratio of {311) <011> and in-plane
anisotropy. Among the samples in Table 2, those indicated as "Inventive steel" in the Remarks
column satiszy the requirements of the steel sheet according to the present invention, and those
indicated as "Comparative steel" in the Remarks column do not satisfy the requirements ofthe
steel sheet according to the present invention. In Table 2, the measurement results that do not
satisfy the requirements of the steel sheet according to the present invention and the
manufacturing conditions that do not satisfy the requirements of the steel sheet manufacturing
method accordingtothe present invention are underlined.
[0140]
As shown in Table 2, in any of the inventive steels B-1, C-1, G-1, H-1, J-1, L-1, P-1,
Q-1, S-1, U-1, W-1, X-1, Y-1, Z-1, AA-1, AB-1 and AD-I, the ratio of the number of carbides
at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds I,
and the Vickers hardness is 150 HV or less. In addition, in any of the inventive steels, the
cross-sectional shrinkage percentage exceeds 40% and the in-plane anisotropy ILI.rl ofthe r
value is 0.2 or less. Thus, they have excellent cold workability. Furthermore, since it was
confirmed that scale scratches were not generated on the steel sheet surface in any of the
inventive steels, these steels can be suitably used for cold working.
- 30-
[0141]
On the other hand, in the Comparative steel A-1, since the AI content is high and the
A3 point decreased, recrystallization during finish hot rolling was inhibited and ILI.rl
deteriorated. Thus, the cold workability is low. In the Comparative steel1-1, the contents of
Mo and Cr are high, recrystallization during finish hot rolling was inhibited, and l~rl
deteriorated. In the comparative steels K-1 and N-1, the content ofS or Mn is high, coarse
MnS was formed in the steel, and the cold workability is low. In the Comparative steel M-1,
the content of Si was high and hardness increased, and thus cold workability is low. Also, in
the Comparative steel M-1, since the A3 point rose, recrystallization during finish hot rolling
was hindered and ILI.rl deteriorated_
[0142]
In the Comparative steel 0-1, C is high, the volume fraction of carbides increased, a
large amount of cracks as the starting point of fractures were generated, and the cross-sectional
shrinkage percentage was low. Thus, the cold workability is low. In the Comparative steel D-
1, the fmish temperature of hot rolling was low and the productivity decreased. In the
Comparative steel F-1, the finish temperature of hot rolling was high, and scale scratches were
generated on the surface of the steel sheet.
[01431
In the Comparative steels R-1 and AC-1, the coiling temperature of hot rolling was
low, the low-temperature transformation structure such as bainite and martensite increased
resulting in brittled steel, and breaks frequently occurred when the hot-rolled coil was
discharged resulting in a decrease in productivity. In the Comparative steels E-1 and T-1, the
coiling temperature of hot rolling was high, thick pearlite with lamellar spacing and needleshaped
coarse carbides with high thermal stability were produced in the hot rolled structure.
Since these carbides remained in the steel sheet even after the two-step type annealing, the
cross-sectional shrinkage percentage was low and thus the cold workability is low.
[0144]
Subsequently, in order to investigate the effect of annealing conditions, steel strips
(slabs) having the ingredient composition shown in Table 1 were heated at 1240°C for 1.8
hours aud then subjected to hot rolling. After completing fmish hot rolling at 890°C, they were
cooled to 520°C at a cooling rate of 45°C/sec on ROT and coiled at 51 ooc to produce a hotrolled
coil with a thickness of3.0 mm. And under the conditions shown in Table 3, a hotrolled
sheet-annealed sample with a thickness of3.0 mm was prepared.
[0145]
For each ofthe samples prepared, the carbide diameter, the ferrite grain diameter, the
Vickers hardness, the ratio of the number of carbides at the ferrite grain boundary relative to
the number of carbides in the ferrite grain, the cross-sectional shrinkage percentage, the X-ray
- 31 -
diffraction intensity ratio of {311} <0 11> and the in-plane anisotropy were determined in the
same manner as the inventive steels and the comparative steels in Table 2. The results are
shown in Table 3.
[Table 3]
1st step annealing 2nd step annealing
Carbide
Retention Retention Retention Retention Cooling diameter
temp. time temp. time rate [~m]
['C] [hr] ['C] [hr] ['C/sec]
A-2 669 27 761 41 7 0.96
B-2 695 47 753 12 30 0.80
C-2 654 19 771 32 20 0.90
D-2 705 51 760 23 13 0.98
E-2 693 50 729 45 34 0.60
F-2 698 22 743 36 14 0.91
G-2 695 55 759 19 7 1.26
ll-2 698 30 742 I 6 0.98
1-2 658 57 754 14 16 0.75
J-2 694 42 811 29 7 1.59
K-2 675 21 746 42 22 0.71
L-2 694 14 781 12 30 0.84
M-2 709 9 779 8 29 0.81
N-2 670 15 738 14 25 0.76
0-2 676 52 732 46 7 0.82
P-2 671 21 769 42 9 1.23
Q:;I 701 2 756 34 25 0.79
R-2 663 52 729 15 17 0.61
S-2 681 56 779 4 7 1.41
T-2 741 47 774 9 25 1.29
U-2 672 12 741 47 25 0.71
W-2 700 45 730 8 28 0.72
X-2 709 46 743 45 8 1.15
Y-2 676 51 777 54 17 0.94
Z-2 668 30 706 23 16 0.57
AA-2 662 14 776 26 19 2.23
AB-2 678 68 785 12 14 1.16
AC-2 637 35 745 48 6 1.19
AD-2 712 50 750 29 15 0.96
Ferrite Grain
grain Vickers boundary Cross~sectional
diameter hardness carbide shrinkage
[HV] No.I grain ratio [%]
[~m] carbide No.
27.1 116 4.8 56.3
15.2 136 3.1 48.7
23.3 112 2.8 58.7
25.4 108 3.7 59.4
13.6 161 1.7 38.6
16.8 120 4.9 54.8
24.2 118 11.4 51.0
13.5 133 0.4 50.9
12.8 133 8.2 53.6
38.1 !51 2.9 32.6
26.5 104 2.7 20.7
22.0 109 4.3 55.9
18.2 183 3.9 38.6
10.9 153 11.2 39.6
14.6 148 4.0 37.4
29.5 108 4.6 59.0
24.9 135 3.2 37.6
10.3 138 4.5 48.0
27.3 126 7.6 49.9
32.6 138 1.5 34.7
17.7 121 3.7 42.1
I 0.3 141 2.1 46.5
22.2 118 4.2 49.6
37.7 141 2.8 29.2
8.5 124 0.5 51.0
54.0 105 9.4 60.2
29.4 116 4.8 45.7
26.2 133 6.8 31.2
19.3 127 2.3 50.9
{311}<011>
X-ray InRplane
diffraction anisotropy
intensity IMI
ratio (11)
4.6 0.38
1.9 0.10
1.8 0.11
1.5 0.06
1.2 0.04
2.5 0.17
2.2 0.15
1.5 0.07
3.8 0.28
1.9 0.10
1.0 0.03
2.4 0.17
2.8 0.18
1.6 0.06
1.2 0.05
1.4 0.05
1.6 0.05
2.0 0.12
1.2 0.04
1.1 0.01
2.6 0.16
2.0 0.11
1.8 0.08
1.8 0.07
1.1 0.04
1.4 0.05
1.4 0.06
1.9 0.10
1.7 0.06
Remarks
Com12.steel
Developed steel
Developed steel
Developed steel
Com]2.steel
Developed steel
Developed steel
Com]2.ste~: is greater than 3.0. In these comparative steels, the in-plane
anisotropy lllrl exceeds 0.2, and thus the cold workability is low. As described above, by
performing analysis by X-ray diffraction on a plane parallel to the plate surface at the 112 plate
thickness portion of the hot-rolled steel sheet, the degree of plastic anisotropy such as the inplane
anisotropy IMI or the quality of cold workability of the hot-rolled steel sheet to be cold
worked can be determined before cold working.
INDUSTRIAL APPLICABILITY
[0155]
- 35 -
As described above, according to the present invention, a steel sheet with excellent
cold workability during forming can be manufactured and provided. The steel sheet of the
present invention is a steel sheet suitable as a material for automotive parts, blades, and other
mechanical parts manufactured through processing steps such as punching, bending, pressing,
etc. Therefore, the present invention has excellent industrial applicability.

CLAIMS
1. A steel sheet having an excellent cold workability during forming, comprising,
in terms of% by mass:
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.0001 to 0.010%,
AI: 0.001 to 0.10%, and
a balance of Fe and inevitable impurities,
wherein (a) a ratio of the number of carbides at a ferrite grain boundary relative to the
number of carbides in the ferrite grain is more than 1,
wherein (b) a diameter of the ferrite grain is 5 flm or more and 50 fLm or less,
wherein (c) an in-plane anisotropy IMI of the r value standardized according to JIS Z
2254 is 0.2 or less,
wherein (d) a Vickers hardness ofthe steel sheet is 100 HV or more and !50 HV or
less, and
wherein (e) a ratio ofX-ray diffraction intensity of the {311} <011> orientation atthe
1/2-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when
a sample with a random orientation distribution of crystal grains in the steel sheet is subjected
to X-ray diffraction is 3.0 or less.
2. The steel sheet with excellent cold workability during forming according to
claim I further comprising, in terms of% by mass, one or a plurality of:
N: 0.0001 to 0.010%,
0: 0.0001 to 0.020%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.10%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%,
Cu: 0.001 to 0.1 0%,
W: 0.001 to 0.10%,
Ta: 0.001 to 0.1 0%,
Ni: 0.001 to 0.10%,
Sn: 0.00 I to 0.050%,
Sb: 0.001 to 0.050%,
As: 0.001 to 0.050%,
Mg: 0.0001 to 0.050%,
Ca: 0.001 to 0.050%,
Y: 0.001 to 0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%, and
Ce: 0.001 to 0.050%.
- 37-
3. A method for producing a steel sheet with excellent cold workability during
forming according to claim 1 or 2, said method comprising:
subjecting a steel strip having an ingredient composition according to claim 1 or 2 to
hot rolling by heating, followed by completing the finish hot rolling at a temperature range of
800°C or higher and 900°C or lower;
coiling said hot-rolled steel sheet at a temperature of 400°C or higher and 550°C or
lower;
pickling said hot-rolled steel sheet, and then subjecting said hot-rolled steel sheet to a
two-step type annealing in which said hot-rolled steel sheet is retained in two temperature
ranges,
. wherein the two-step type annealing comprises
(i) subjecting said hot-rolled steel sheet to a first step annealing performed
by retaining said hot-rolled steel at a temperature range of 650°C or higher and 720°C or lower
for 3 hours or longer and 60 hours or shorter, and then a second step annealing performed by
retaining the hot-rolled steel at a temperature range of725°C or higher and 790°C or lower for
3 hours or longer and 50 hours or shorter, and thereafter
(ii) cooling said hot-rolled steel sheet to 650°C or lower at a cooling rate of
1 °C/hour or more and 30°C/hour or less.
4. The method for producing a steel sheet according to claim 3, wherein the steel
sheet has a cross-sectional shrinkage percentage of 40% or more.