[0001]The present invention relates to a steel sheet and a method of producing the same. More
particularly, the present invention relates to a high-strength steel sheet having excellent stretch
10 formability, and a method of producing the same.
BACKGROUND
[0002]
In order to improve the stretch formability of a DP steel (dual-phase steel mainly composed
15 of ferrite and martensite) having a tensile strength of 550 MPa or higher and 1,100 MPa or lower,
it is desirable to integrate the crystal orientation of b.c.c. (body-centered cubic lattice) to -fibers.
In addition, it is necessary to reduce the integration of the orientation to other than -fibers as
much as possible. In DP steels, a high strength is achieved by utilizing martensite structure;
however, martensite is sometimes accumulated in a specific orientation. This is caused by the
20 formation of an austenite crystallographic texture and, specifically, the formation of an austenite
crystallographic texture in the orientations referred to as “copper orientation” and “brass
orientation” leads to the formation of a crystallographic texture in martensite that is generated
when austenite is cooled. The information of this martensite crystallographic texture is expressed
by ODF (crystal orientation distribution function) (2  45); however, since the martensite
25 crystallographic texture exists on -fibers, it is difficult to recognize the difference thereof from
parent-phase ferrite.
[0003]
Numerous inventions relating to DP steels and high-strength steel sheets have been
disclosed to date; however, only a handful of them disclose a technology relating to an
30 improvement of the stretch formability (see, for example, PTLs 1 to 4).
[0004]
PTL 1 discloses a technology relating to a high-tensile-strength hot-rolled steel sheet having
high formability with excellent stretch flange formability and fatigue characteristics as well as
good stretch formability and shape fixability, in which a steel slab containing C: 0.01 to 0.10%
35 by weight, Si: 0.50 to 1.50% by weight, Mn: 0.50 to 2.50% by weight, P: 0.05% by weight or
less, S: 0.005% by weight or less, and Ti: 0.005 to 0.03% by weight is retained in a temperature
2
range of 900 to 1,300C, subsequently continuously hot rolled at a rolling reduction ratio of
lower than 20% in the final stand and a rolling termination temperature of 870 to 980C, cooled
at a cooling rate of 50 to 200C/sec after the completion of the rolling, and then wound in a coil
form in a temperature range of 300 to 650C, whereby a structure which is composed of a ferrite
phase having a volume ratio of 70 to 97% and the 5 remainder being a low-temperature
transformed phase mainly constituted by a bainite phase is formed, and the in-plane anisotropy
r of r value is controlled to be 0.2 or less. It is noted here, however, that PTL1 does not offer a
technology for ensuring the formability for the case of a steel structure containing martensite
structure beneficial for improvement of the strength.
10 [0005]
PTL 2 discloses a technology relating to a high-strength cold-rolled steel sheet having a
small in-plane anisotropy of elongation and a tensile strength (TS) of 440 MPa or higher with
excellent press formability, in which the steel sheet has a composition containing, by mass %, C:
0.030 to 0.20%, Si: 1.5% or less, Mn: 1.0 to 2.5%, P: 0.005 to 0.1%, S: 0.01% or less, Al: 0.005
15 to 1.5%, N: 0.01%, and the balance of Fe and unavoidable impurities, and the steel sheet
contains, by area ratio with respect to the whole steel sheet structure: 85% to 99% of a ferrite
phase as parent phase; and 1% to 15% of a second phase containing a martensite phase, the area
ratio of the martensite phase with respect to the whole steel sheet structure is 1% to 13%, and the
average crystal orientation density (I) of  fibers in a range of   25 to 35, which is expressed
20 by an ODF (crystal orientation distribution function), is 2.0 to 4.0 in the crystallographic texture
of a sheet plane at a 1/4-thickness position of the steel sheet. In this technology, the area ratio of
the martensite structure is reduced to decrease the in-plane anisotropy and, therefore, a high
strength and a high ductility, which are characteristics of a DP steel, cannot be obtained. It can
be understood also from this disclosed technology that it is necessary to modify the martensite
25 structure in order to improve the stretch formability while maintaining the characteristics of a
conventional DP steel.
[0006]
PTL 3 discloses a high-strength hot-dip galvanized steel sheet having excellent formability
along with a TS of 780 MPa or higher, excellent elongation El and a TS  EL value of 18,000 or
30 more, wherein the steel sheet has a component composition that contains, by mass %, C: 0.03 to
0.15%, Si: 0.8 to 2.5%, Mn: 1.0 to 3.0%, P: 0.001 to 0.05%, S: 0.0001 to 0.01%, Al: 0.001 to
0.1%, N: 0.0005 to 0.01%, Cr: 0.1 to 2.0%, and the balance of Fe and unavoidable impurities,
and has a microstructure that contains, by area ratio, 50% or more of a ferrite phase and 10% or
more of a martensite phase. In PTL 3, only a technology of improving the bulging height by
35 providing the steel sheet surface with a plated coating film and a post-treatment coating film is
3
disclosed, and no technology is presented with regard to the isotropy of the shape after forming,
which is an important indicator of stretch formability.
[0007]
PTL 4 discloses a high-strength steel sheet with excellent workability, which has a tensile
strength of 590 MPa or higher and in which 5 uniform elongation and hole expansion are improved
simultaneously, wherein the steel sheet contains, by mass %, C: 0.04 to 0.10%, Mn: 0.5 to 2.6%,
and Si: 0.8 to 2.0%, the ratio (C/Si) of the C content to the Si content being 0.04 or more and less
than 0.10; the content of Al, P, S and N being restricted; and the steel sheet comprises a
metallographic structure composed of, by volume ratio, 90 to 95% of ferrite and 5 to 10% of
10 tempered martensite. This disclosed technology is nothing more than a means of tempering the
martensite structure and reducing the area ratio of tempered martensite for an improvement the
workability; therefore, PTL 4 still has room for improvement in terms of improving the stretch
formability.
In addition to the above, for example, PTLs 5 to 7 each disclose a technology relating to a
15 high-strength steel sheet; however, the stretch formability is not examined at all in these PTLs.
[CITATION LIST]
[PATENT LITERATURE]
[0008]
20 [PTL 1] JP 2000-297349 A
[PTL 2] JP 2009-132981 A
[PTL 3] JP 2010-236027 A
[PTL 4] JP 2011-032543 A
[PTL 5] JP 2016-130357 A
25 [PTL 6] JP 2016-130355 A
[PTL 7] JP 2015-193897 A
SUMMARY
[TECHNICAL PROBLEM]
30 [0009]
In view of the above-described circumstances, an object of the present invention is to
provide a steel sheet having a high strength and excellent stretch formability, and a method of
producing the same.
35 [SOLUTION TO PROBLEM]
[0010]
4
The present inventors conducted intensively studies on the techniques for solving the
above-described problems, and examined the changes in orientation in detail so as to distinguish
the development of martensite crystallographic texture. As a result, it was revealed that, by
reducing the integration of the orientation referred to as “(252)<2-11>“, not only the formation
of martensite crystallographic texture 5 can be inhibited (the orientation integration degree of
martensite can be randomized), but also the stretch formability can be improved (the anisotropy
can be reduced). It was also found that this orientation appears after the transformation of
austenite in the copper orientation and the brass orientation into martensite and is not visually
recognizable in conventional ODF (2  45).
10 [0011]
Further, the present inventors conducted various studies to discover that it is difficult to
produce a steel sheet with limited integration of the above-described orientation even if the
hot-rolling conditions, the annealing conditions and the like are simply and individually devised,
and that such a steel sheet can be produced only by achieving optimization in a so-called
15 consistent process of the hot rolling and annealing step and the like, thereby completing the
present invention.
[0012]
The gist of the present invention is as follows.
[0013]
20 (1) A steel sheet, having a chemical composition comprising, by mass %:
C: 0.05 to 0.20%;
Si: 0.01 to 1.30%;
Mn: 1.00 to 3.00%;
P: 0.0001 to 0.0200%;
25 S: 0.0001 to 0.0200%;
Al: 0.001 to 1.000%;
N: 0.0001 to 0.0200%;
Co: 0 to 0.5000%;
Ni: 0 to 0.5000%;
30 Mo: 0 to 0.5000%;
Cr: 0 to 1.0000%;
O: 0 to 0.0200%;
Ti: 0 to 0.5000%;
B: 0 to 0.0100%;
35 Nb: 0 to 0.5000%;
V: 0 to 0.5000%;
5
Cu: 0 to 0.5000%;
W: 0 to 0.1000%;
Ta: 0 to 0.1000%;
Sn: 0 to 0.0500%;
5 Sb: 0 to 0.0500%;
As: 0 to 0.0500%;
Mg: 0 to 0.0500%;
Ca: 0 to 0.0500%;
Y: 0 to 0.0500%;
10 Zr: 0 to 0.0500%;
La: 0 to 0.0500%;
Ce: 0 to 0.0500%; and
a balance of Fe and impurities,
wherein
15 the steel sheet comprises, by area ratio:
a total of ferrite and bainite: 10.0 to 90.0%;
a total of martensite and tempered martensite: 5.0 to 80.0%; and
a total of pearlite and retained austenite: 0 to 15.0%,
integration degree of (111)<112> orientation of ferrite is 3.0 or higher; and
20 integration degree of (252)<2-11> orientation of martensite and tempered martensite is 5.0
or lower.
(2) The steel sheet according to (1), containing one or more of:
Co: 0.0001 to 0.5000%;
Ni: 0.0001 to 0.5000%;
25 Mo: 0.0001 to 0.5000%;
Cr: 0.0001 to 1.0000%;
O: 0.0001 to 0.0200%;
Ti: 0.0001 to 0.5000%;
B: 0.0001 to 0.0100%;
30 Nb: 0.0001 to 0.5000%;
V: 0.0001 to 0.5000%;
Cu: 0.0001 to 0.5000%;
W: 0.0001 to 0.1000%;
Ta: 0.0001 to 0.1000%;
35 Sn: 0.0001 to 0.0500%;
Sb: 0.0001 to 0.0500%;
6
As: 0.0001 to 0.0500%;
Mg: 0.0001 to 0.0500%;
Ca: 0.0001 to 0.0500%;
Y: 0.0001 to 0.0500%;
5 Zr: 0.0001 to 0.0500%;
La: 0.0001 to 0.0500%; and
Ce: 0.0001 to 0.0500%.
(3) A method of producing a steel sheet, the method comprising:
a casting step of continuously casting a molten steel having the chemical composition
10 according to (1) or (2) to form a slab, wherein 5 to 40% rolling reduction is performed at a
temperature of 800C to lower than 1,200C in a period after the continuous casting and before
cooling to room temperature;
a hot rolling step which includes hot rolling the slab and in which a finishing temperature of
the hot rolling is 650 to 950C;
15 a step of coiling the thus obtained hot-rolled steel sheet at a coiling temperature of 400 to
700C;
a step of retaining the thus coiled hot-rolled steel sheet as is wherein the thus coiled
hot-rolled steel sheet is not cooled to room temperature before retaining and is retained in a
temperature range of (coiling start temperature + 20C to 100C) for 5 to 300 minutes,;
20 a cold rolling step of cold rolling the hot-rolled steel sheet at a rolling reduction ratio of
10.0 to 90.0%; and
an annealing step of annealing the thus obtained cold-rolled steel sheet in a temperature
range of 700 to 900C.
25 [ADVANTAGEOUS EFFECTS OF INVENTION]
[0014]
According to the present invention, a steel sheet having a high strength and excellent stretch
formability, and a method of producing the same can be provided.
30 [BRIEF DESCRIPTION OF DRAWING]
[0015]
FIG. 1 is a graph showing the effects of the integration degree of the (111)<112>
orientation of ferrite and the integration degree of the (252)<2-11> orientation of martensite and
tempered martensite on the stretch formability of the DP steels used in Examples 1 and 2.
35
DESCRIPTION OF EMBODIMENTS
7
[0016]
Embodiments of the present invention will now be described. It is noted here, however, the
following descriptions are merely intended for exemplification of the embodiments of the
present invention, and the present invention is not limited to the below-described embodiments.
5 [0017]

The steel sheet according to one embodiment of the present invention has a chemical
composition comprising, by mass %:
C: 0.05 to 0.20%;
10 Si: 0.01 to 1.30%;
Mn: 1.00 to 3.00%;
P: 0.0001 to 0.0200%;
S: 0.0001 to 0.0200%;
Al: 0.001 to 1.000%;
15 N: 0.0001 to 0.0200%;
Co: 0 to 0.5000%;
Ni: 0 to 0.5000%;
Mo: 0 to 0.5000%;
Cr: 0 to 1.0000%;
20 O: 0 to 0.0200%;
Ti: 0 to 0.5000%;
B: 0 to 0.0100%;
Nb: 0 to 0.5000%;
V: 0 to 0.5000%;
25 Cu: 0 to 0.5000%;
W: 0 to 0.1000%;
Ta: 0 to 0.1000%;
Sn: 0 to 0.0500%;
Sb: 0 to 0.0500%;
30 As: 0 to 0.0500%;
Mg: 0 to 0.0500%;
Ca: 0 to 0.0500%;
Y: 0 to 0.0500%;
Zr: 0 to 0.0500%;
35 La: 0 to 0.0500%;
Ce: 0 to 0.0500%; and
8
a balance of Fe and impurities,
wherein
the steel sheet comprises, by area ratio:
a total of ferrite and bainite: 10.0 to 90.0%;
a 5 total of martensite and tempered martensite: 5.0 to 80.0%; and
a total of pearlite and retained austenite: 0 to 15.0%,
the integration degree of the (111)<112> orientation of ferrite is 3.0 or higher; and
the integration degree of (252)<2-11> orientation of martensite and tempered martensite is
5.0 or lower.
10 [0018]
First, the reasons for restricting the chemical components of the steel sheet according to one
embodiment of the present invention will be described. Hereinafter, “%” used for each
component means “% by mass”.
[0019]
15 (C: 0.05 to 0.20%)
C is an element which inexpensively increases the tensile strength, and is an extremely
important factor for controlling the orientation integration degree of ferrite and bainite or
martensite and tempered martensite. When the C content is less than 0.05%, retained austenite
cannot be stabilized at the time of hot rolling and coiling, and the orientation integration degree
20 of martensite cannot be randomized. Accordingly, a lower limit value is set at 0.05% or more.
The C content may be 0.06% or more, 0.07% or more, or 0.08% or more. When the C content is
more than 0.20%, the stretch formability is deteriorated since such a C content causes not only a
reduction in elongation but also a reduction of the orientation integration degree of ferrite.
Accordingly, an upper limit value is set at 0.20% or less. The C content may be 0.18% or less,
25 0.16% or less, or 0.15% or less.
[0020]
(Si: 0.01 to 1.30%)
Si is an element which acts as a deoxidizer and affects the form of carbides and heat-treated
retained austenite. In order to satisfy both wear resistance and stretch formability, it is effective
30 to improve the strength by reducing the volume ratio of carbides existing in a steel component
and utilizing retained austenite. When the Si content is less than 0.01%, the generation of
carbides is not inhibited and a large amount of carbides thus exists in the steel, as a result of
which the stretch formability is deteriorated. Accordingly, a lower limit value is set at 0.01% or
more. The Si content may be 0.05% or more, 0.10% or more, or 0.30% or more. Further, the Si
35 content of more than 1.30% causes an increase in the steel strength and embrittlement of the steel
component, consequently deteriorating the stretch formability. Accordingly, an upper limit value
9
is set at 1.30% or less. The Si content may be 1.20% or less, 1.10% or less, 1.00% or less, or
0.90% or less.
[0021]
(Mn: 1.00 to 3.00%)
Mn, which is a factor affecting ferrite 5 transformation of steel, is an element effective for
improving the strength. When the Mn content is less than 1.00%, martensitic transformation
cannot be facilitated in the cooling process of cold-rolled sheet annealing, and this causes a
reduction in the strength. Accordingly, a lower limit value is set at 1.00% or more. The Mn
content may be 1.10% or more, 1.30% or more, or 1.50% or more. When the Mn content is more
10 than 3.00%, the stretch formability is deteriorated since ferrite transformation and bainite
transformation in cold-rolled sheet annealing are inhibited. Accordingly, an upper limit value is
set at 3.00% or less. The Mn content may be 2.80% or less, 2.50% or less, or 2.20% or less.
[0022]
(P: 0.0001 to 0.0200%)
15 P is an element which is strongly segregated at ferrite grain boundaries to facilitate
embrittlement of the grain boundaries. The lower the P content, the more preferred it is. When
the P content is less than 0.0001%, a long time is required for refining to increase the purity, and
this leads to a significant increase in the cost. Accordingly, a lower limit value is set at 0.0001%
or more. The P content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more. When
20 the P content is more than 0.0200%, the stretch formability is deteriorated due to grain boundary
embrittlement. Accordingly, an upper limit value is set at 0.0200% or less. The P content may be
0.0180% or less, 0.0150% or less, or 0.0120% or less.
[0023]
(S: 0.0001 to 0.0200%)
25 S is an element which generates non-metallic inclusions such as MnS in steel and causes a
reduction in the ductility of a steel component. The lower the S content, the more preferred it is.
When the S content is less than 0.0001%, a long time is required for refining to increase the
purity, and this leads to a significant increase in the cost. Accordingly, a lower limit value is set
at 0.0001% or more. The S content may be 0.0005% or more, 0.0010% or more, or 0.0020% or
30 more. When the S content is more than 0.0200%, the stretch formability is deteriorated since
such an S content causes cracking that originate from non-metallic inclusions during cold
forming. Accordingly, an upper limit value is set at 0.0200% or less. The S content may be
0.0180% or less, 0.0150% or less, or 0.0120% or less.
[0024]
35 (Al: 0.001 to 1.000%)
10
Al is an element which acts as a deoxidizer of steel and stabilizes ferrite, and Al is added as
required. When the Al content is less than 0.001%, the effect of the addition is not sufficiently
obtained. Accordingly, a lower limit value is set at 0.001% or more. The Al content may be
0.005% or more, 0.010% or more, or 0.020% or more. When the Al content is more than 1.000%,
the strength of the steel sheet is reduced 5 since ferrite transformation and bainite transformation
are excessively facilitated in the cooling process of cold-rolled sheet annealing. Accordingly, an
upper limit value is set at 1.000% or less. The Al content may be 0.950% or less, 0.900% or less,
or 0.800% or less.
[0025]
10 (N: 0.0001 to 0.0200%)
N is an element which forms coarse nitrides in the steel sheet and deteriorates the
workability of the steel sheet. N is also an element which causes generation of blow-holes in
welding. An N content of less than 0.0001% leads to a significant increase in the production cost.
Accordingly, a lower limit value is set at 0.0001% or more. The N content may be 0.0005% or
15 more, 0.0010% or more, or 0.0020% or more. When the N content is more than 0.0200%,
deterioration of the stretch formability and generation of blow-holes become prominent.
Accordingly, an upper limit value is set at 0.0200% or less. The N content may be 0.0180% or
less, 0.0160% or less, or 0.0120% or less.
[0026]
20 The basic chemical composition of the steel sheet according to one embodiment of the
present invention is as described above. The steel sheet may further contain the following
elements as required. The steel sheet may contain the following elements in place of a part of the
balance of Fe.
[0027]
25 (Co: 0 to 0.5000%)
Co is an element effective for controlling the form of carbides and improving the strength,
and it is added as required. When the Co content is less than 0.0001%, the effects of the addition
are not obtained. Accordingly, a lower limit value is preferably set at 0.0001% or more. The Co
content may be 0.0002% or more, 0.0010% or more, or 0.0100% or more. When the Co content
30 is more than 0.5000%, fine Co carbide precipitates in a large amount, and this causes an increase
in the strength and a reduction in the ductility of the steel material, as a result of which the cold
workability and the stretch formability may be deteriorated. Accordingly, an upper limit value is
set at 0.5000% or less. The Co content may be 0.4500% or less, 0.4000% or less, or 0.3000% or
less.
35 [0028]
(Ni: 0 to 0.5000%)
11
Ni is a reinforcing element and is effective for improving the hardenability. In addition, Ni
may be added since it improves the wettability and facilitates an alloying reaction. When the Ni
content is less than 0.0001%, these effects are not obtained. Accordingly, a lower limit value is
preferably set at 0.0001% or more. The Ni content may be 0.0002% or more, 0.0010% or more,
or 0.0100% 5 or more. When the Ni content is more than 0.5000%, the productivity in the
production and in hot rolling may be adversely affected, or the stretch formability may be
deteriorated. Accordingly, an upper limit value is set at 0.5000% or less. The Ni content may be
0.4500% or less, 0.4000% or less, or 0.3000% or less.
[0029]
10 (Mo: 0 to 0.5000%)
Mo is an element effective for improving the strength of the steel sheet. Further, Mo is an
element which has an effect of inhibiting ferrite transformation that occurs during a heat
treatment performed in a continuous annealing equipment or continuous hot-dip galvanizing
equipment. When the Mo content is less than 0.0001%, these effects are not obtained.
15 Accordingly, a lower limit value is preferably set at 0.0001% or more. The Mo content may be
0.0002% or more, 0.0010% or more, or 0.0100% or more. When the Mo content is more than
0.5000%, the formability, particularly the stretch formability may be deteriorated since not only
ferrite transformation and bainite transformation are inhibited but also martensitic transformation
is facilitated during cold-rolled sheet annealing. Accordingly, an upper limit value is set at
20 0.5000% or less. The Mo content may be 0.4500% or less, 0.4000% or less, or 0.3000% or less.
[0030]
(Cr: 0 to 1.0000%)
Similarly to Mn, Cr is an element which inhibits pearlite transformation and is effective for
improving the steel strength, and Cr is added as required. When the Cr content is less than
25 0.0001%, the effects of the addition are not obtained. Accordingly, a lower limit value is
preferably set at 0.0001% or more. The Cr content may be 0.0002% or more, 0.0010% or more,
or 0.0100% or more. When the Cr content is more than 1.0000%, the stretch formability may be
deteriorated since the stability of austenite is markedly increased and a large amount of retained
austenite thus exists after cold-rolled sheet annealing. Accordingly, an upper limit value is set at
30 1.0000% or less. The Cr content may be 0.9000% or less, 0.8000% or less, or 0.7000% or less.
[0031]
(O: 0 to 0.0200%)
O forms oxides and deteriorates the workability; therefore, the amount thereof to be added
needs to be kept small. Particularly, the oxides often exist in the form of inclusions and, when
35 such oxides exist on a punched end surface or a cut surface, notch-like defects and coarse
dimples are formed on the end surface, as a result of which stress concentration is induced during
12
stretch forming and severe working, and the workability is significantly deteriorated with such
defects and dimples serving as the origin of crack formation. However, an O content of less than
0.0001% is not economically preferred since it leads to an excessively high cost. Accordingly, a
lower limit value is preferably set at 0.0001% or more. The O content may be 0.0005% or more,
0.0010% or more, or 0.0020% 5 or more. Meanwhile, when the O content is more than 0.0200%,
the above-described tendency of workability deterioration is pronounced. Accordingly, an upper
limit value is set at 0.0200% or less. The O content may be 0.0180% or less, 0.0150% or less, or
0.0100% or less.
[0032]
10 (Ti: 0 to 0.5000%)
Ti is a reinforcing element. Ti contributes to an increase in the strength of the steel sheet
through strengthening by precipitates, fine-grain strengthening by the inhibition of the growth of
ferrite crystal grains, and dislocation strengthening by the inhibition of recrystallization. When
the Ti content is less than 0.0001%, these effects are not obtained. Accordingly, a lower limit
15 value is preferably set at 0.0001% or more. Ti content may be 0.0002% or more, 0.0010% or
more, or 0.0100% or more. When the Ti content is more than 0.5000%, the formability,
particularly the stretch formability may be deteriorated due to an increased precipitation of
carbonitrides. Accordingly, an upper limit value is set at 0.5000% or less. The Ti content may be
0.4500% or less, 0.4000% or less, or 0.3000% or less.
20 [0033]
(B: 0 to 0.0100%)
B is an element which inhibits the generation of ferrite and pearlite from austenite in a
cooling process and facilitates the generation of a low-temperature transformed structure of
bainite, martensite or the like. Further, B is an element beneficial for improving the steel strength,
25 and it is added as required. When the B content is less than 0.0001%, the effect of improving the
strength or the wear resistance by the addition is not sufficiently obtained. Moreover, not only
the most careful attention must be paid when performing an analysis to identify a B content of
less than 0.0001%, but also such a B content may be below the detection limit depending on the
analysis equipment. Accordingly, a lower limit value is preferably set at 0.0001% or more. The
30 B content may be 0.0003% or more, 0.0005% or more, or 0.0010% or more. When the B content
is more than 0.0100%, coarse B oxide may be generated in the steel, and the stretch formability
may be deteriorated with the B oxide serving as the origin of void generation during cold
forming. Accordingly, an upper limit value is set at 0.0100% or less. The B content may be
0.0080% or less, 0.0060% or less, or 0.0050% or less.
35 [0034]
(Nb: 0 to 0.5000%)
13
Similarly to Ti, Nb is an element effective for controlling the form of carbides and, since an
addition thereof leads to structural refinement, Nb is also an element effective for improving the
toughness. When the Nb content is less than 0.0001%, these effects are not obtained.
Accordingly, a lower limit value is preferably set at 0.0001% or more. The Nb content may be
0.0002% or more, 0.0010% or more, or 5 0.0100% or more. When the Nb content is more than
0.5000%, fine and hard Nb carbide precipitates in a large amount, and this causes an increase in
the strength and a marked reduction in the ductility of the steel material, as a result of which the
cold workability and the stretch formability may be deteriorated. Accordingly, an upper limit
value is set at 0.5000% or less. The Nb content may be 0.4500% or less, 0.4000% or less, or
10 0.3000% or less.
[0035]
(V: 0 to 0.5000%)
V is a reinforcing element. V contributes to an increase in the strength of the steel sheet
through strengthening by precipitates, fine-grain strengthening by the inhibition of the growth of
15 ferrite crystal grains, and dislocation strengthening by the inhibition of recrystallization. When
the V content is less than 0.0001%, these effects are not obtained. Accordingly, a lower limit
value is preferably set at 0.0001% or more. The V content may be 0.0002% or more, 0.0010% or
more, or 0.0100% or more. When the V content is more than 0.5000%, the formability,
particularly the stretch formability is deteriorated due to an increased precipitation of
20 carbonitrides. Accordingly, an upper limit value is set at 0.5000% or less. The V content may be
0.4500% or less, 0.4000% or less, or 0.3000% or less.
[0036]
(Cu: 0 to 0.5000%)
Cu is an element effective for improving the strength of the steel sheet. When the Cu
25 content is less than 0.0001%, this effect is not obtained. Accordingly, a lower limit value is
preferably set at 0.0001% or more. The Cu content may be 0.0002% or more, 0.0010% or more,
or 0.0100% or more. When the Cu content is more than 0.5000%, the steel material is embrittled
during hot rolling, making it impossible to perform hot rolling. In addition, the stretch
formability may be deteriorated due to a marked increase in the steel strength. Accordingly, an
30 upper limit value is set at 0.5000% or less. The Cu content may be 0.4500% or less, 0.4000% or
less, or 0.3000% or less.
[0037]
(W: 0 to 0.1000%)
W is an extremely important element not only because it is effective for improving the
35 strength of the steel sheet, but also because W-containing precipitates and crystals act as
hydrogen trapping sites. When the W content is less than 0.0001%, these effects are not obtained.
14
Accordingly, a lower limit value is preferably set at 0.0001% or more. The W content may be
0.0002% or more, 0.0010% or more, or 0.0050% or more. When the W content is more than
0.1000%, the workability, particularly the stretch formability may be deteriorated. Accordingly,
an upper limit value is set at 0.1000% or less. The W content may be 0.0800% or less, 0.0600%
5 or less, or 0.0500% or less.
[0038]
(Ta: 0 to 0.1000%)
Similarly to Nb, V and W, Ta is an element effective for controlling the form of carbides
and improving the strength, and Ta is added as required. When the Ta content is less than
10 0.0001%, the effects of the addition are not obtained. Accordingly, a lower limit value is
preferably set at 0.0001% or more. The Ta content may be 0.0002% or more, 0.0010% or more,
or 0.0050% or more. When the Ta content is more than 0.1000%, fine Ta carbide is precipitated
in a large amount, and this causes an increase in the strength and a reduction in the ductility of
the steel sheet, as a result of which the bending resistance and the stretch formability may be
15 deteriorated. Accordingly, an upper limit value is set at 0.1000% or less. The Ta content may be
0.0800% or less, 0.0600% or less, or 0.0500% or less.
[0039]
(Sn: 0 to 0.0500%)
Sn is an element which is incorporated into steel when scrap is used as a raw material, and
20 the lower the Sn content, the more preferred it is. An Sn content of less than 0.0001%, however,
leads to an increase in the refining cost. Accordingly, a lower limit value is preferably set at
0.0001% or more. The Sn content may be 0.0002% or more, 0.0010% or more, or 0.0050% or
more. When the Sn content is more than 0.0500%, the stretch formability may be deteriorated
due to embrittlement of ferrite. Accordingly, an upper limit value is set at 0.0500% or less. The
25 Sn content may be 0.0400% or less, 0.0300% or less, or 0.0200% or less.
[0040]
(Sb: 0 to 0.0500%)
Similarly to Sn, Sb is an element which is incorporated when scrap is used as a steel raw
material. Sb is strongly segregated at grain boundaries and causes embrittlement of the grain
30 boundaries and a reduction of the ductility; therefore, the lower the Sb content, the more
preferred it is, and the Sb content may be 0%. An Sb content of less than 0.0001%, however,
leads to an increase in the refining cost. Accordingly, a lower limit value is preferably set at
0.0001% or more. The Sb content may be 0.0002% or more, 0.0010% or more, or 0.0050% or
more. When the Sb content is more than 0.0500%, the stretch formability may be deteriorated.
35 Accordingly, an upper limit value is set at 0.0500% or less. The Sb content may be 0.0400% or
less, 0.0300% or less, or 0.0200% or less.
15
[0041]
(As: 0 to 0.0500%)
Similarly to Sn and Sb, As is an element which is incorporated when scrap is used as a steel
raw material, and is strongly segregated at grain boundaries. The lower the As content, the more
preferred it is. An As content of 5 less than 0.0001%, however, leads to an increase in the refining
cost. Accordingly, a lower limit value is preferably set at 0.0001% or more. The As content may
be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the As content is more than
0.0500%, the stretch formability is deteriorated. Accordingly, an upper limit value is set at
0.0500% or less. The As content may be 0.0400% or less, 0.0300% or less, or 0.0200% or less.
10 [0042]
(Mg: 0 to 0.0500%)
Mg is an element which can control the form of sulfides when added in a trace amount, and
it is added as required. When the Mg content is less than 0.0001%, this effect is not obtained.
Accordingly, a lower limit value is preferably set at 0.0001% or more. The Mg content may be
15 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the Mg content is more than
0.0500%, the stretch formability may be deteriorated due to the formation of coarse inclusions.
Accordingly, an upper limit value is set at 0.0500% or less. The Mg content may be 0.0400% or
less, 0.0300% or less, or 0.0200% or less.
[0043]
20 (Ca: 0 to 0.0500%)
Ca is useful as a deoxidizing element and also exerts an effect in controlling the form of
sulfides. When the Ca content is less than 0.0001%, the effects of Ca are not sufficiently
obtained. Accordingly, a lower limit value is preferably set at 0.0001% or more. The Ca content
may be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the Ca content is more
25 than 0.0500%, the workability, particularly the stretch formability may be deteriorated.
Accordingly, an upper limit value is set at 0.0500% or less. The Ca content may be 0.0400% or
less, 0.0300% or less, or 0.0200% or less.
[0044]
(Y: 0 to 0.0500%)
30 Similarly to Mg and Ca, Y is an element which can control the form of sulfides when added
in a trace amount, and it is added as required. When the Y content is less than 0.0001%, this
effect is not obtained. Accordingly, a lower limit value is preferably set at 0.0001% or more. The
Y content may be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the Y content
is more than 0.0500%, the stretch formability may be deteriorated due to the formation of coarse
35 Y oxide. Accordingly, an upper limit value is set at 0.0500% or less. The Y content may be
0.0400% or less, 0.0300% or less, or 0.0200% or less.
16
[0045]
(Zr: 0 to 0.0500%)
Similarly to Mg, Ca and Y, Zr is an element which can control the form of sulfides when
added in a trace amount, and it is added as required. When the Zr content is less than 0.0001%,
this effect is not obtained. Accordingly, a 5 lower limit value is preferably set at 0.0001% or more.
The Zr content may be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the Zr
content is more than 0.0500%, the stretch formability may be deteriorated due to the formation
of coarse Zr oxide. Accordingly, an upper limit value is set at 0.0500% or less. The Zr content
may be 0.0400% or less, 0.0300% or less, or 0.0200% or less.
10 [0046]
(La: 0 to 0.0500%)
La is an element which is effective for controlling the form of sulfides when added in a
trace amount, and it is added as required. When the La content is less than 0.0001%, this effect is
not obtained. Accordingly, a lower limit value is preferably set at 0.0001% or more. The La
15 content may be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the La content
is more than 0.0500%, the stretch formability may be deteriorated due to the formation of La
oxide. Accordingly, an upper limit value is set at 0.0500% or less. The La content may be
0.0400% or less, 0.0300% or less, or 0.0200% or less.
[0047]
20 (Ce: 0 to 0.0500%)
Similar to La, Ce is an element which can control the form of sulfides when added in a trace
amount, and it is added as required. When the Ce content is less than 0.0001%, this effect is not
obtained. Accordingly, a lower limit value is preferably set at 0.0001% or more. The Ce content
may be 0.0002% or more, 0.0010% or more, or 0.0050% or more. When the Ce content is more
25 than 0.0500%, the stretch formability may be deteriorated due to the formation of Ce oxide.
Accordingly, an upper limit value is set at 0.0500% or less. The Ce content may be 0.0400% or
less, 0.0300% or less, or 0.0200% or less.
[0048]
In the steel sheet according to one embodiment of the present invention, the remainder other
30 than the above-described components is composed of Fe and impurities. The term “impurities”
used herein includes components which are incorporated due to various factors of the production
process during the industrial production of a steel sheet, such as raw materials including ore,
scrap and the like, and are not intentionally added to the steel sheet according to one embodiment
of the present invention (so-called unavoidable impurities). The term “impurities” also includes
35 elements other than the above-described components, which elements are contained in the steel
17
sheet according to one embodiment of the present invention at a level that the actions and effects
unique to the respective elements do not affect the properties of the steel sheet.
[0049]
Next, the characteristic features of the structure and properties of the steel sheet according
5 to one embodiment of the present invention will be described.
[0050]
(Total of Ferrite and Bainite: 10.0 to 90.0%)
The total area ratio of ferrite and bainite affects the elongation of steel, and the workability
is improved as the area ratio increases. A total area ratio of lower than 10.0% requires an
10 advanced control in the production; therefore, the yield may be reduced and the stretch
formability may be deteriorated. Accordingly, a lower limit value is set at 10.0% or higher. The
total area ratio of ferrite and bainite may be 20.0% or higher, 30.0% or higher, or 35.0% or
higher. When the total area ratio is higher than 90%, the strength may be reduced. Accordingly,
an upper limit value is set at 90.0% or lower. The total area ratio of ferrite and bainite may be
15 85.0% or lower, 80.0% or lower, or 75.0% or lower.
[0051]
(Total of Martensite and Tempered Martensite: 5.0 to 80.0%)
The total area ratio of martensite and tempered martensite affects the steel strength, and the
tensile strength is increased as the area ratio increases. At lower than 5.0%, the area ratio of
20 martensite and tempered martensite is not sufficient, and a target tensile strength of 550 MPa or
higher may not be achieved. Accordingly, a lower limit value is set at 5.0% or higher. The total
area ratio of martensite and tempered martensite may be 10.0% or higher, 15.0% or higher, or
20.0% or higher. When the total area ratio is higher than 80.0%, the tensile strength may exceed
1,100 MPa to cause deterioration of the strength-ductility balance and the stretch formability.
25 Accordingly, an upper limit value is set at 80.0% or lower. The total area ratio of martensite and
tempered martensite may be 70.0% or lower, 60.0% or lower, or 55.0% or lower.
[0052]
(Total of Pearlite and Retained Austenite: 0 to 15.0%)
Pearlite and retained austenite of the balance are structural factors that deteriorate the local
30 ductility of steel; therefore, the lower the content thereof, the more preferred it is. The total area
ratio of pearlite and retained austenite may be 0%; however, a total area ratio of lower than 1.0%
may require an advanced control in the production. From the standpoint of preventing a
reduction in the yield, the total area ratio of pearlite and retained austenite can be set at 1.0% or
higher. The total area ratio of pearlite and retained austenite may be 2.0% or higher, 3.0% or
35 higher, or 5.0% or higher. When this total area ratio is higher than 15.0%, the stretch formability
18
may be deteriorated. Accordingly, an upper limit value is set at 15.0% or lower. The total area
ratio of pearlite and retained austenite may be 13.0% or lower, 11.0% or lower, or 9.0% or lower.
[0053]
(Integration Degree of (111)<112> Orientation of Ferrite: 3.0 or Higher)
The integration degree of the (5 111)<112> orientation of ferrite is a factor that affects
isotropic deformation of steel, i.e. stretch formability, and the higher this integration degree, the
better the stretch formability. When the integration degree is lower than 3.0, good stretch
formability cannot be obtained. Accordingly, a lower limit value is set at 3.0 or higher. The
integration degree is preferably 4.0 or higher, or 5.0 or higher. An upper limit value of this
10 integration degree is not particularly limited, and it may be 10.0 or lower, 8.0 or lower, or 7.0 or
lower.
[0054]
(Integration Degree of (252)<2-11> Orientation of Martensite and Tempered Martensite: 5.0 or
lower)
15 The integration degree of the (252)<2-11> orientation of martensite and tempered
martensite in total is a factor that inhibits isotropic deformation of steel, i.e. a factor that affects
the stretch formability, and the lower this integration degree, the better the stretch formability.
When the integration degree is higher than 5.0, the stretch formability is deteriorated.
Accordingly, an upper limit value is set at 5.0 or lower. The integration degree is preferably 4.0
20 or lower, or 3.0 or lower. A lower limit value of this integration degree is not particularly limited,
and it may be 0.1 or higher, 0.2 or higher, or 0.3 or higher.
[0055]
(Sheet Thickness)
The thickness of the steel sheet is a factor that affects the rigidity of a steel member after
25 forming, and the greater the sheet thickness, the higher the rigidity of the member. When the
sheet thickness is less than 0.2 mm, not only the rigidity is reduced but also the stretch
formability is deteriorated due to the effects of unavoidable non-metallic inclusions existing in
the steel material; therefore, the sheet thickness is preferably 0.2 mm or greater. However, when
the sheet thickness is greater than 3.0 mm, the forming load is increased at the time of stretch
30 forming, and this causes wear of a die and deterioration of the productivity; therefore, the sheet
thickness is preferably 3.0 mm or less.
[0056]
Next, method of observing and measuring the above-prescribed structures will be described.
[0057]
35 (Method of Evaluating Total Area Ratio of Ferrite and Bainite)
19
The area ratio of ferrite and bainite is determined by observing a portion in a range of 1/8 to
3/8 of the sheet thickness that is centered at the 1/4-thickness position on an electron channeling
contrast image under a field emission-scanning electron microscope (FE-SEM). Electron
channeling contrast imaging is a technique for detecting misorientation in crystal grains as a
difference in contrast and, on 5 the thus obtained image, polygonal ferrite is observed as a part
having a uniform contrast within a structure judged as ferrite, not pearlite, bainite, martensite or
retained austenite. Bainite is a collection of lath-like crystal grains inside of which iron-based
carbides with a major axis of 20 nm or longer are not contained, or inside of which iron-based
carbides with a major axis of 20 nm or longer are contained and the carbides belong to a single
10 variant, i.e. a group of iron-based carbides extending in the direction. The phrase “group of
iron-based carbides extending in the same direction” used herein refers to a group of iron-based
carbides in which a difference in the extension direction is within 5. Bainite surrounded by
grain boundaries having an orientation difference of 15 or larger is counted as a single bainite
grain. The total area ratio of ferrite and bainite is determined by an image analysis method for
15 each of eight viewing fields on a 35 m  25 m electron channeling contrast image, and an
average value thereof is defined as the total area ratio of ferrite and bainite.
[0058]
(Method of Evaluating Total Area Ratio of Martensite and Tempered Martensite)
The total area ratio of martensite and tempered martensite is also determined from the
20 above-described image taken by electron channeling contrast imaging. The structures of
martensite and tempered martensite are less likely to be etched than ferrite and thus exist as
protrusions on the structure observation surface. It is noted here that tempered martensite is a
collection of lath-like crystal grains, inside of which iron-based carbides with a major axis of 20
nm or longer are contained and the carbides belong to plural variants, i.e. plural groups of
25 iron-based carbides extending in different directions. Further, retained austenite also exist as
protrusions on the structure observation surface. Therefore, the total area ratio of martensite and
tempered martensite can be accurately measured by subtracting the area ratio of retained
austenite that is determined by the below-described procedures from the area ratio of the
protrusions that is determined by the above-described procedures.
30 [0059]
(Method of Evaluating Total Area Ratio of Pearlite and Retained Austenite)
The area ratio of retained austenite can be determined by a measurement using an X-ray. In
other words, a portion of a sample from a sheet surface to the 1/4-depth position in the sheet
thickness direction is removed by mechanical polishing and chemical polishing. Subsequently,
35 the fraction of retained austenite structure is calculated from the integrated intensity ratios of the
(200) and (211) diffraction peaks of the bcc phase and the (200), (220) and (311) diffraction
20
peaks of the fcc phase, which are obtained by using MoK radiation as a characteristic X-ray on
the polished sample, and the thus calculated value is defined as the area ratio of retained
austenite. Further, the area ratio of pearlite is determined from an image taken by the
above-described electron channeling contrast imaging. Pearlite is a structure in which plate-like
5 carbide and ferrite are layered.
[0060]
(Method of Evaluating Integration Degree of (111)<112> Orientation of Ferrite)
The orientation integration degree of ferrite is measured using an EBSD (Electron Back
Scattering Diffraction) apparatus, and can be measured by either an EBSP (Electron Back
10 Scattering Pattern) method or an ECP (Electron Channeling Pattern) method. The orientation
integration degree of ferrite may be determined from a three-dimensional crystallographic
texture calculated by a vector method based on a 110 pole figure, or from a three-dimensional
crystallographic texture calculated by a series expansion method using plural (preferably three or
more) pole figures among 110, 100, 211, and 310 pole figures. In the measurement by
15 EBSD, the crystal orientation data at the same position as that of the above-described electron
channeling contrast image are obtained with the STEP interval being set at 0.05 m. The
integration degree of the (111)<112> orientation is determined from the crystal orientation data
corresponding to ferrite that are obtained by this procedure for eight viewing fields.
[0061]
20 (Method of Evaluating Integration Degree of (252)<2-11> Orientation of Martensite and
Tempered Martensite in Total)
The orientation integration degree of martensite and tempered martensite is also determined
by EBSD. The crystal orientation data collected for the evaluation of the orientation integration
degree of ferrite also contain the crystal orientation data of martensite and tempered martensite.
25 In the same manner as in the case of ferrite, the integration degree of the (252)<2-11> orientation
is determined from the crystal orientation data of martensite and tempered martensite on an
electron channeling contrast image.
[0062]
(Mechanical Properties)
30 According to the steel sheet of one embodiment of the present invention, the stretch
formability can be improved while achieving a high tensile strength and an excellent
strength-ductility balance, specifically a tensile strength of 550 to 1,100 MPa and a total
elongation of 10.0% or more. The tensile strength is preferably 700 MPa or higher, more
preferably 800 MPa or higher.
35 [0063]

21
A method of producing the steel sheet according to one embodiment of the present
invention is characterized by coherent management of the hot rolling, cold rolling and annealing
conditions with the use of materials in the above-described component ranges. One example of a
method of producing a steel sheet will now be described; however, a method of producing the
steel sheet according to the present i 5 nvention is not restricted to the below-described mode.
The method of producing the steel sheet according to one embodiment of the present
invention is characterized by including:
a casting step of continuously casting a molten steel having the chemical composition as
that described above for the steel sheet to form a slab (steel piece), wherein 5 to 40% rolling
10 reduction is performed at a temperature of 800C to lower than 1,200C in a period after the
continuous casting and before cooling to room temperature;
a hot rolling step which includes hot rolling the slab and in which a finishing temperature of
the hot rolling is 650 to 950C;
a step of coiling the thus obtained hot-rolled steel sheet at a coiling temperature of 400 to
15 700C;
a step of retaining the thus coiled hot-rolled steel sheet as is in a temperature range of
(coiling start temperature + 20C to 100C) for 5 to 300 minutes, without cooling the steel sheet
to room temperature;
a cold rolling step of cold rolling the hot-rolled steel sheet at a rolling reduction ratio of
20 10.0 to 90.0%; and
an annealing step of annealing the thus obtained cold-rolled steel sheet in a temperature
range of 700 to 900C.
These steps will now each be described in detail.
[0064]
25 (Casting Step)
In the method of producing the steel sheet according to one embodiment of the present
invention, first, a slab is formed by continuously casting a molten steel having the same chemical
composition as that described above for the steel sheet, and 5 to 40% rolling reduction is
subsequently performed at a temperature of 800C to lower than 1,200C in a period after the
30 continuous casting and before cooling of the slab to room temperature, whereby the uniformity
of microsegregation-concentrated parts in the slab is improved (specifically,
element-concentrated parts can be finely dispersed in the steel material to reduce the difference
in concentration of the element-concentrated parts). When the rolling reduction ratio is lower
than 5%, segregation is not eliminated, as a result of which the orientation integration degree of
35 ferrite and bainite is reduced and the stretch formability is deteriorated. By improving the
uniformity of the element-concentrated parts in the slab (e.g., improving the uniformity of
22
Mn-concentrated parts), non-recrystallized ferrite is prevented from remaining in the
element-concentrated parts after cold rolling and annealing, and the orientation is integrated to
the (111) plane of ferrite, so that a stretch-formed section is made likely to isotropically expand.
In addition, in the below-described post-coiling retention step, austenite is likely to be generated
in the hot-rolled sheet. Accordingly, a lower limit 5 value of the rolling reduction ratio is set at 5%
or higher, and the rolling reduction ratio may be 6% or higher, 8% or higher, or 10% or higher.
When the rolling reduction ratio is higher than 40%, an increase in the size of the equipment is
required, and this leads to a heavy equipment investment and an increase in the cost. Further,
since the growth directions of solidified structures are aligned, the orientation integration degree
10 of ferrite and bainite is reduced after cold-rolled sheet annealing due to the effect of the resulting
crystallographic texture of the solidified structures, as a result of which the stretch formability is
deteriorated. Accordingly, an upper limit value is set at 40% or lower, and it may be 38% or
lower, 35% or lower, or 30% or lower.
[0065]
15 (Hot Rolling Step)
In the present method, hot rolling is subsequently performed on the cast slab. This hot
rolling step can be performed by reheating and hot rolling the cast slab directly, or after once
cooling the cast slab. When reheating is performed, the heating temperature of the slab is
generally 1,100C or higher. An upper limit value thereof is not particularly defined, and may be,
20 for example, 1,250C or lower.
[0066]
(Rough Rolling)
In the present method, for example, rough rolling may be optionally performed on the cast
slab before finish-rolling so as to adjust the resulting sheet thickness and the like. The conditions
25 of this rough rolling are not particularly restricted as long as the desired sheet bar dimensions can
be ensured.
[0067]
(Finish Rolling)
Subsequently, finish rolling is performed on the thus obtained slab, or the slab that has been
30 additionally rough rolled as required and, in this process, the finishing temperature (hot-rolling
finishing temperature of the hot rolling) is controlled in a range of 650 to 950C. The hot-rolling
finishing temperature is a factor that affects the control of the crystallographic texture of prior
austenite grains. When the hot-rolling finishing temperature is lower than 650C, the rolled
crystallographic texture of austenite is developed, and anisotropy is induced in the steel material
35 properties. Accordingly, a lower limit value is set at 650C or higher, and the hot-rolling
finishing temperature may be 680C or higher, or 700C or higher. When the hot-rolling
23
finishing temperature is higher than 950C, since the material is maintained at a high temperature
prior to the rolling, abnormal grain growth of austenite occurs, making it difficult to obtain an
isotropic crystallographic texture. Accordingly, an upper limit value is set at 950C or lower, and
the hot-rolling finishing temperature may be 930C or lower, or 900C or lower.
5 [0068]
(Coiling Step)
After the hot rolling step, the thus obtained hot-rolled steel sheet is coiled at a coiling
temperature of 400 to 700C in the subsequent coiling step. The coiling temperature is an
important factor for controlling ferrite and bainite that transform from austenite in the structural
10 change of the hot-rolled sheet. When the coiling temperature is lower than 400C, even with the
below-described post-coiling heating treatment, austenite existing in the hot-rolled sheet after the
coiling cannot be transformed into bainite, and a target hot-rolled structure cannot be obtained.
In addition, the stretch formability is deteriorated as a result. Accordingly, a lower limit value is
set at 400C or higher, and the coiling temperature may be 420C or higher, or 450C or higher.
15 When the coiling temperature is higher than 700C, the transformation from austenite to ferrite is
excessively facilitated during the coiling of the hot-rolled sheet, carbon is concentrated in
austenite, and pearlite transformation is proceeded when applicating the below-described
post-coiling heating treatment, therefore a target hot-rolled structure cannot be obtained.
Accordingly, an upper limit value is set at 700C or lower, and the coiling temperature may be
20 680C or lower, or 650C or lower.
[0069]
(Retention Step)
Next, the thus coiled hot-rolled steel sheet is, without being cooled to room temperature
before retaining, retained as is in a temperature range of (coiling start temperature + 20C to
25 100C) over a period of 5 to 300 minutes. Such an increase in the temperature and the retention
in the temperature range of (coiling start temperature + 20C to 100C) are extremely important
control factors in the present invention. During the process of cooling the hot-rolled sheet and
terminating the cooling to the coiling temperature after the hot finish rolling, ferrite
transformation or bainite transformation proceeds, and carbon is gradually concentrated to the
30 residual austenite. This reaction also proceeds after the hot-rolled sheet is wound into a coil form,
and the Mn concentration on the austenite side of austenite-B.C.C. interface in the hot-rolled
sheet structure is reduced by once increasing the temperature after the ferrite transformation or
bainite transformation; therefore, the austenite-B.C.C. interface is allowed to move and
eventually, stable retained austenite can be obtained in the state of a hot-rolled sheet even at
35 room temperature. As described above, in the method of producing the steel sheet according to
one embodiment of the present invention, the uniformity of the element-concentrated parts in the
24
slab is improved by controlling the rolling reduction conditions of the slab in the casting step. By
combining this with the temperature retention conditions in the retention step, austenite can be
generated and allowed to remain in the resulting hot-rolled sheet more appropriately. This
retained austenite stabilized in the state of a hot-rolled sheet remains to exist even after cold
rolling. The retained austenite generated in this hot rolled she 5 et by heat treatment is mixed with
austenite generated from the crystallographic texture of ferrite in the K-S relation during cold
rolling and annealing, and the crystallographic texture of austenite in the cold-rolled sheet
annealing is thereby randomized, as a result of which the integration degree of the (252)<2-11>
orientation of martensite in the final product can be reduced. By setting the retention temperature
10 at (coiling start temperature + 20C) or higher, not only the movement of the interface in the
transformation from untransformed austenite to bainite and the growth of bainite structure can be
promoted, but also the concentration of carbon to the retained austenite can be facilitated.
Further, by controlling the retention temperature at (coiling start temperature + 100C) or lower,
internal oxidation can be inhibited. When the retention time is shorter than 5 minutes, the effects
15 of the present invention cannot be obtained since the stabilization of austenite due to the progress
of bainite transformation is insufficient. Accordingly, a lower limit value is set at 5 minutes or
longer, and the retention time may be 15 minutes or longer, or 30 minutes or longer. When the
retention time is longer than 300 minutes, oxygen is supplied from the steel strip surface to the
inside, and internal oxides are thereby formed in the hot-rolled sheet. The “internal oxides” are
20 oxides formed along the grain boundaries, and serve as the origin of cracks when they remain
after cold rolling and annealing, causing deterioration of the stretch formability. Accordingly, an
upper limit value is set at 300 minutes or shorter, and the retention time may be 250 minutes or
shorter, or 200 minutes or shorter.
[0070]
25 (Cold Rolling and Annealing Step)
Lastly, the thus obtained hot-rolled steel sheet is, for example, pickled as required, and
subsequently cold rolled at a rolling reduction ratio of 10.0 to 90.0% and annealed at a
temperature of 700 to 900C, whereby the steel sheet according to one embodiment of the
present invention is obtained. In the method of producing the steel sheet according to one
30 embodiment of the present invention, the retained austenite generated in the hot-rolled sheet in
the above-described casting step and retention step and the austenite newly generated by the cold
rolling and annealing both remain after the cold rolling and annealing. In other words, different
orientations of austenite remain and co-exist. In this manner, by combining the rolling reduction
conditions in the casting step, the temperature retention conditions in coiling and the cold-rolling
35 and annealing conditions and thereby allowing different orientations of austenite to remain, the
integration degree of the (252)<2-11> orientation of martensite and tempered martensite in the
25
steel sheet to be eventually obtained can be more appropriately and easily reduced. Preferred
embodiments of cold rolling, annealing and plating treatment are described below in detail. The
descriptions below are, however, merely examples of preferred embodiments of cold rolling,
annealing and plating treatment, and should not restrict a method of producing the steel sheet by
5 any means.
[0071]
(Pickling)
First, prior to the cold rolling, the coiled hot-rolled steel sheet is uncoiled and pickled. By
performing this pickling, oxide scales on the surface of the hot-rolled steel sheet can be removed
10 to improve the chemical conversion and plating properties of the resulting cold-rolled steel sheet.
The pickling may be performed once or plural separate times.
[0072]
(Cold-Rolling Reduction Ratio)
The cold-rolling reduction ratio affects the recrystallization behavior of ferrite during the
15 cold rolling and annealing. In addition, the cold-rolling reduction has an effect in that the crystal
orientation of the retained austenite existing in the hot-rolled sheet is rotated by the cold rolling,
and the crystal orientation of the austenite generated by the cold rolling and annealing is
randomized. When the cold-rolling reduction ratio is lower than 10.0%, the orientation
integration degree of ferrite is reduced, and the stretch formability is deteriorated. Accordingly, a
20 lower limit value is set at 10.0% or higher, and the cold-rolling reduction ratio may be 15.0% or
higher. When the cold-rolling reduction ratio is higher than 90.0%, although ferrite is easily
recrystallized, austenite generated in the hot-rolled sheet undergoes strain-induced
transformation, and the orientation integration degree of martensite and tempered martensite is
increased; therefore, the stretch formability is deteriorated. Accordingly, an upper limit value is
25 set at 90.0% or lower, and the cold-rolling reduction ratio may be 75.0% or lower.
[0073]
(Cold-Rolled Sheet Annealing)
(Heating Rate)
When the cold-rolled steel sheet is passed through a continuous annealing line or a plating
30 line, the heating rate is not particularly restricted; however, since the productivity may be largely
deteriorated at a heating rate of lower than 0.5C/sec, the heating rate is preferably 0.5C/sec or
higher. On the other hand, a heating rate of higher than 100C/sec involves an excessively large
equipment investment; therefore, the heating rate is preferably 100C/sec or lower.
[0074]
35 (Annealing Temperature)
26
The annealing temperature is a factor that affects the recrystallization behavior of ferrite.
The annealing temperature is also a control factor that affects the behavior of austenite
generation and is an extremely important for controlling the strength-ductility balance of steel.
When the annealing temperature is lower than 700C, the amount of generated austenite is small,
and undissolved carbides exist even after the retention 5 in the cold rolling and annealing. Further,
the transformation from austenite to pearlite is facilitated by the presence of undissolved
carbides; therefore, in the resulting cold-rolled and annealed structure, the ratio of martensite
structure is reduced while the ratio of pearlite structure is increased. In addition, since
non-recrystallized ferrite remains, the stretch formability is deteriorated. Accordingly, a lower
10 limit value is set at 700C or higher, and the annealing temperature may be 750C or higher.
When the annealing temperature is higher than 900C, since the amount of austenite generated
during the retention at a constant temperature in the annealing is increased, the orientation
integration degree of ferrite and bainite is reduced in the resulting cold-rolled and annealed
structure, as a result of which the stretch formability is deteriorated. Accordingly, an upper limit
15 value is set at 900C or lower, and the annealing temperature may be 850C or lower.
[0075]
(Retention Time)
The steel sheet is supplied to a continuous annealing line to perform annealing with heating
at the annealing temperature. In this process, the retention time is preferably 10 to 600 seconds.
20 When the retention time is shorter than 10 seconds, the fraction of austenite at the annealing
temperature is insufficient and/or the carbides existing prior to the annealing are not sufficiently
dissolved, as a result of which prescribed structure and properties may not be obtained. A
retention time of longer than 600 seconds presents no problem in terms of properties; however,
since it requires a long equipment line, an upper limit is substantially about 600 seconds.
25 [0076]
(Average Cooling Rate)
After the above-described annealing, cooling is preferably performed from 750C to 550C
at an average cooling rate of 100.0C/sec or lower. A lower limit value of the average cooling
rate is not particularly restricted and may be, for example, 2.5C/sec. The reason for setting the
30 lower limit value of the average cooling rate at 2.5C/sec is to inhibit the occurrence of ferrite
transformation in the base steel sheet and thereby prevent the base steel sheet from being
softened. When the average cooling rate is lower than 2.5C/sec, the strength may be reduced.
The average cooling rate is more preferably 5.0C/sec or higher, still more preferably 10.0C/sec
or higher, yet still more preferably 20.0C/sec or higher. At a temperature of higher than 750C,
35 the cooling rate is not restricted since ferrite transformation is unlikely to occur. At a temperature
of lower than 550C, the cooling rate is also not restricted since a low-temperature transformed
27
structure is obtained. When the cooling is performed at a rate of higher than 100.0C/sec, a
low-temperature transformed structure is generated in the surface layer as well, and this causes a
variation in hardness; therefore, the cooling is performed at a rate of preferably 100.0C/sec or
lower, more preferably 80.0C/sec or lower, still more preferably 60.0C/sec or lower.
5 [0077]
(Cooling Stop Temperature)
The above-described cooling is stopped at a temperature of 25C to 550C (cooling stop
temperature). Subsequently, when this cooling stop temperature is lower than (plating bath
temperature - 40C), the steel sheet may be reheated and retained in a temperature range of
10 350C to 550C. When the cooling is performed in the above-described temperature range,
martensite is generated from untransformed austenite during the cooling. By reheating the steel
sheet thereafter, martensite is tempered, and precipitation of carbides as well as recovery and
rearrangement of dislocations take place in the hard phase, as a result of which the hydrogen
embrittlement resistance is improved. The reason why the lower limit of the cooling stop
15 temperature is set at 25C is not only because excessive cooling requires a significant equipment
investment, but also because the effects of the cooling are saturated.
[0078]
(Retention Temperature)
After the reheating but before immersion in a plating bath, the steel sheet may be retained in
20 a temperature range of 350 to 550C. The retention in this temperature range not only contributes
to tempering of martensite, but also eliminates temperature variation of the sheet in the width
direction and improves the post-plating outer appearance. It is noted here that, when the cooling
stop temperature is 350C to 550C, the steel sheet may be retained as is without reheating.
[0079]
25 (Retention Time)
The duration of the retention is desirably set at 10 seconds to 600 seconds so as to obtain
the effects of the retention.
[0080]
(Tempering)
30 In a series of annealing operations, the cold-rolled sheet, or the cold-rolled sheet on which a
plating treatment has been performed, may be started to be reheated after being cooled to room
temperature or in the middle of being cooled to room temperature (provided that the sheet
temperature is not higher than the martensitic transformation start temperature (Ms)), and then
retained in a temperature range of 150C to 400C for 2 seconds or longer. According to this step,
35 by tempering martensite generated during the post-reheating cooling into tempered martensite,
the hydrogen embrittlement resistance can be improved. When the tempering step is performed,
28
with a retention temperature of lower than 150C or a retention time of shorter than 2 seconds,
martensite is not sufficiently tempered, and satisfactory changes thus may not be attained in
terms of microstructure and mechanical properties. On the other hand, a retention temperature of
higher than 400C causes a reduction of the dislocation density in tempered martensite, as a
result of which the tensile strength may be deteriorated. 5 Therefore, when tempering is performed,
it is preferred to retain the steel sheet in a temperature range of 150C to 400C for 2 seconds or
longer. The tempering may be performed inside a continuous annealing equipment, or may be
performed using a separate off-line equipment after the continuous annealing. In this process, the
tempering time varies depending on the tempering temperature. In other words, the lower the
10 tempering temperature, the longer is the tempering time, and the higher the tempering
temperature, the shorter is the tempering time.
[0081]
(Plating)
During or after the annealing step, as required, hot-dip galvanization may be performed on
15 the cold-rolled steel sheet by heating or cooling the cold-rolled steel sheet to a temperature of
(galvanizing bath temperature - 40)C to (galvanizing bath temperature + 50)C. By this hot-dip
galvanization step, a hot-dip galvanized layer is formed on at least one surface, preferably both
surfaces of the cold-rolled steel sheet. In this case, the corrosion resistance of the cold-rolled
steel sheet is improved, which is preferred. Even when hot-dip galvanization is performed, the
20 hydrogen embrittlement resistance of the cold-rolled steel sheet can be maintained sufficiently.
[0082]
For a plating treatment, for example, the Sendzimir method in which “after degreasing and
pickling, a steel sheet is heated in a non-oxidizing atmosphere, annealed in a reducing
atmosphere containing H2 and N2, subsequently cooled to the vicinity of the temperature of a
25 plating bath, and then immersed in the plating bath”, a total reduction furnace method in which
“after the atmosphere during annealing is adjusted and a steel sheet surface is oxidized first, the
steel sheet surface is reduced and thereby cleaned before being plated, and subsequently
immersed in a plating bath”, or a flux method in which “after degreasing and pickling of a steel
sheet, the steel sheet is flux-treated with ammonium chloride and subsequently immersed in a
30 plating bath” may be employed, and the effects of the present invention can be exerted under any
of these treatment conditions.
[0083]
(Plating Bath Temperature)
The plating bath temperature is preferably 450 to 490C. When the plating bath temperature
35 is lower than 450C, the viscosity of the plating bath is excessively increased and this makes it
difficult to control the thickness of the plated layer, as a result of which the outer appearance of
29
the resulting hot-dip galvanized steel sheet may be deteriorated. On the other hand, when the
plating bath temperature is higher than 490C, a large amount of fume is generated, and this can
make it difficult to safely perform the plating operations. The plating bath temperature is more
preferably 455C or higher, but it is more preferably 480C or lower.
5 [0084]
(Composition of Plating Bath)
As for the composition of the plating bath, the plating bath is preferably mainly composed
of Zn and has an effective Al amount (a value obtained by subtracting a total Fe content from a
total Al content in the plating bath) of 0.050 to 0.250% by mass. When the effective Al amount
10 in the plating bath is less than 0.050% by mass, the plating adhesion may be deteriorated due to
excessive diffusion of Fe into the plated layer. On the other hand, when the effective Al amount
in the plating bath is greater than 0.250% by mass, Al-based oxides that inhibit the movement of
Fe atoms and Zn atoms are generated at the interface between the steel sheet and the plated layer,
as a result of which the plating adhesion may be deteriorated. The effective Al amount in the
15 plating bath is more preferably 0.065% by mass or greater, but it is more preferably 0.180% by
mass or less.
[0085]
(Steel Sheet Temperature at Immersion in Plating Bath)
The plating bath immersion sheet temperature (the temperature of the steel sheet at the time
20 of being immersed in a hot-dip galvanizing bath) is preferably in a range of 40C lower than the
hot-dip galvanizing bath temperature (“hot-dip galvanizing bath temperature - 40C”) to 50C
higher than the hot-dip galvanizing bath temperature (“hot-dip galvanizing bath temperature +
50C”). A plating bath immersion sheet temperature of lower than [hot-dip galvanizing bath
temperature - 40C] is not desirable since this may lead to deterioration of the plated outer
25 appearance due to a large heat loss during the immersion in the plating bath and partial
solidification of molten zinc. When the sheet temperature prior to the immersion is lower than
[hot-dip galvanizing bath temperature - 40C], the steel sheet may be further heated prior to the
immersion in the plating bath by an arbitrary method so as to control the sheet temperature to be
[hot-dip galvanizing bath temperature - 40C] or higher, and the steel sheet may be immersed
30 into the plating bath thereafter. Further, when the plating bath immersion sheet temperature is
higher than [hot-dip galvanizing bath temperature + 50C], an operational problem is induced in
association with an increase in the plating bath temperature.
[0086]
(Plating Pretreatment)
30
In order to further improve the plating adhesion, prior to the annealing in a continuous
hot-dip galvanization line, the base steel sheet may be plated with one or more of Ni, Cu, Co,
and Fe.
[0087]
5 (Plating Post-treatment)
On the surface of the hot-dip galvanized steel sheet or alloyed hot-dip galvanized steel sheet,
upper-layer plating and various treatments, such as a chromate treatment, a phosphate treatment,
a lubricity improvement treatment and a weldability improvement treatment, may also be
performed for the purpose of improving the coating properties and the weldability.
10 [0088]
(Skin Pass Rolling)
In addition, skin pass rolling may be performed for the purpose of improving the ductility
through correction of the steel sheet shape and introduction of mobile dislocations. In the skin
pass rolling after the heat treatment, the rolling reduction ratio is preferably in a range of 0.1 to
15 1.5%. A lower limit thereof is set at 0.1% since a rolling reduction ratio of lower than 0.1% has a
small effect and is difficult to control. The productivity is markedly deteriorated when the rolling
reduction ratio is higher than 1.5%; therefore, an upper limit thereof is set at 1.5%. The skin pass
may be performed in-line or off-line. Further, the skin pass of the target rolling reduction ratio
may be performed at once, or may be performed in several separate operations.
20 [0089]
The steel sheet according to the present invention can be obtained by the above-described
production method. In the above, a mode in which the uniformity of
microsegregation-concentrated parts in a slab is improved by controlling the rolling reduction
ratio at 5% or higher in the casting step is described; however, it is also possible to improve the
25 uniformity of microsegregation-concentrated parts by, for example, controlling the temperature
of the slab in the casting step.
EXAMPLES
[0090]
30 Examples of the present invention will now be described; however, the present invention is
not restricted to the conditions of the below-described Examples. The present invention can
adopt a variety of conditions without departing from the gist of the present invention, as long as
the object of the present invention is achieved.
[0091]
35 [Example 1]
31
Steels having the respective chemical compositions shown in Table 1 were each melted and
continuously cast to produce a slab (steel piece), and the thus obtained slab was subsequently
roll-reduced by 6% at a temperature of 800C to lower than 1,200C in a period after the
continuous casting and before cooling of the slab to room temperature, whereby a slab in which
the uniformity of microsegregation-5 concentrated parts was improved (the concentration
difference of element-concentrated parts was reduced) was produced. This slab was inserted to a
furnace heated to 1,220C and retained therein for 60 minutes to perform a homogenization
treatment, after which the slab was taken out to the atmosphere and then hot-rolled to obtain a
steel sheet having a thickness of 2.8 mm. In this hot rolling, the finish-rolling termination
10 temperature (finishing temperature) was 920C and, after a lapse of 1.5 seconds from the
termination of the finish rolling, the steel sheet was water-cooled to a coiling temperature of
610C at a rate of 28C/sec, and then retained at 660C for 1 hour to perform a reheating
treatment. Subsequently, oxide scales on the thus obtained hot-rolled steel sheet were removed
by pickling, and the hot-rolled steel sheet was cold-rolled at a rolling reduction ratio of 50.0% to
15 attain a thickness of 1.4 mm. Further, this cold-rolled steel sheet was heated to 790C at a rate of
8.0C/sec and retained at 790C for 105 seconds, after which the steel sheet was cooled to 480C
at an average cooling rate of 4.0C/sec and subsequently retained at 460C for 12 seconds to
perform cold-rolled sheet annealing. Thereafter, the steel sheet was further skin-pass rolled such
that the resulting steel strip had an elongation of 0.3%. Table 2 shows the results of evaluating
20 the properties of each steel sheet on which the above-described thermo-mechanical treatments
were performed. It is noted here that the remainder other than the components shown in Table 1
was composed of Fe and impurities. The chemical composition analyzed for a sample collected
from each of the thus produced steel sheets was the same as the chemical composition of the
corresponding steel shown in Table 1.
25 [0092]
(Method of Evaluating Tensile Properties)
The tensile strength (TS) and the total elongation (El) were measured by performing a
tensile test in accordance with JIS Z2241(2011) for a JIS No. 5 test piece that was collected such
that the longitudinal direction of the test piece was aligned parallel to the direction perpendicular
30 to the rolling direction of a steel strip.
[0093]
(Method of Evaluating Stretch Formability)
The stretch formability was evaluated by performing the following spherical stretch forming
test.
35 • Sample drawing width: 200 mm  200 mm
• Mold: a spherical punch of 60 mm in radius, beaded die
32
• Pressing load: 60 t
• Bulging rate: 30 mm/min
• Lubrication: application of anti-rust oil
For a steel sheet on which a stretching process was performed to a height of 25 mm under
the above-described conditions, the bulg 5 ing height of the outer surface of the spherically
stretched steel sheet was measured along a circumferential shape at a position of 25 mm away
from the center axis of the spherical punch using a laser or LED noncontact-type displacement
meter. A passing grade (○) was given when the difference between the maximum bulging height
and the minimum bulging height was 3 mm or less, while a failing grade () was given when this
10 difference in height was larger than 3 mm.
[0094]
A steel sheet was evaluated to have a high strength and excellent stretch formability when
the tensile strength was 550 to 1,100 MPa and the evaluation of the stretch formability was “○”.
P190781WO
33
[0095]
[Table 1-1]
Table 1-1
No. C Si Mn P S Al N Co Ni Mo Cr O Ti B Nb
A 0.10 0.23 2.73 0.0012 0.0010 0.104 0.0164 0.0380 0.0404 0.4190 0.1036 0.0042 0.0274 0.0011 0.0265
B 0.16 1.11 2.14 0.0015 0.0033 0.072 0.0014 0.0417 0.1149 0.3876 0.0467 0.0079 0.3734 0.0009 0.4319
C 0.15 0.50 2.01 0.0030 0.0172 0.121 0.0018 0.0456 0.4100 0.0412 0.7489 0.0023 0.0450 0.0012 0.0651
D 0.13 0.72 2.31 0.0163 0.0020 0.041 0.0011 0.0493 0.0543 0.0460 0.1023 0.0009 0.4233 0.0009 0.3895
E 0.08 0.45 2.01 0.0091 0.0032 0.032 0.0029 - - - - - - - -
F 0.11 0.10 1.39 0.0021 0.0021 0.078 0.0015 0.0496 0.0236 0.0592 0.0804 0.0017 0.0795 0.0075 0.3169
G 0.09 0.45 2.58 0.0103 0.0038 0.021 0.0023 - - - - - - - -
H 0.19 0.59 1.87 0.0125 0.0019 0.076 0.0014 0.0708 0.0334 0.3107 0.0420 0.0020 0.0489 0.0061 0.2022
I 0.13 0.38 1.90 0.0020 0.0129 0.141 0.0012 0.3751 0.0333 0.0393 0.2383 0.0012 0.0901 0.0009 0.1212
J 0.18 0.36 2.51 0.0025 0.0020 0.086 0.0074 0.0542 0.0876 0.0391 0.1055 0.0169 0.3238 0.0041 0.0356
K 0.12 0.78 1.38 0.0013 0.0030 0.060 0.0015 - - - - - - - -
L 0.19 0.81 2.80 0.0011 0.0019 0.127 0.0031 - - - - - - - -
M 0.06 0.95 1.66 0.0034 0.0153 0.769 0.0016 0.3032 0.0235 0.0232 0.1936 0.0025 0.0671 0.0010 0.0377
N 0.07 0.11 2.24 0.0118 0.0043 0.613 0.0041 - - - - - - - -
O 0.14 1.15 1.04 0.0009 0.0074 0.250 0.0046 0.1116 0.2078 0.1238 0.1324 0.0126 0.0531 0.0082 0.0702
P 0.07 1.07 2.38 0.0017 0.0015 0.813 0.0018 0.1800 0.0794 0.2034 0.0962 0.0034 0.0307 0.0011 0.0408
Q 0.09 1.23 1.71 0.0152 0.0053 0.098 0.0155 - - - - - - - -
R 0.17 0.09 2.92 0.0014 0.0009 0.610 0.0020 0.4155 0.0440 0.0852 0.6283 0.0017 0.0482 0.0007 0.0394
S 0.04 0.33 1.09 0.0016 0.0021 0.087 0.0020 - - - - - - - -
T 0.21 0.63 2.90 0.0040 0.0023 0.054 0.0164 - - - - - - - -
U 0.13 1.35 2.57 0.0034 0.0137 0.056 0.0173 - - - - - - - -
V 0.09 1.07 0.93 0.0039 0.0144 0.061 0.0014 - - - - - - - -
W 0.11 0.26 3.06 0.0163 0.0024 0.116 0.0020 - - - - - - - -
X 0.13 0.15 1.16 0.0206 0.0016 0.777 0.0008 - - - - - - - -
Bold and underlined values are outside the scope of the present invention.
5
[0096]
[Table 1-2]
P190781WO
34
Table 1-2
No. C Si Mn P S Al N Co Ni Mo Cr O Ti B Nb
Y 0.10 0.22 1.54 0.0012 0.0205 0.144 0.0020 0.3324 0.0506 0.1718 0.6617 0.0030 0.3446 0.0010 0.0503
Z 0.12 0.21 2.69 0.0020 0.0091 1.038 0.0018 - - - - - - - -
AA 0.06 0.69 1.80 0.0134 0.0053 0.052 0.0207 0.0402 0.0386 0.0615 0.0687 0.0010 0.4016 0.0016 0.0262
AB 0.14 0.69 1.36 0.0018 0.0014 0.106 0.0013 0.5102 0.1532 0.1044 0.8161 0.0019 0.0692 0.0024 0.0214
AC 0.18 0.84 1.89 0.0027 0.0073 0.301 0.0012 0.3917 0.5168 0.0595 0.0457 0.0018 0.1047 0.0007 0.0424
AD 0.16 1.14 2.19 0.0153 0.0010 0.116 0.0024 0.0518 0.0906 0.5153 0.3512 0.0016 0.0359 0.0012 0.0276
AE 0.11 0.40 2.86 0.0140 0.0031 0.836 0.0165 0.3616 0.0627 0.0223 1.0306 0.0042 0.0454 0.0011 0.0511
AF 0.06 0.89 2.87 0.0017 0.0016 0.166 0.0028 0.3095 0.0520 0.0345 0.1101 0.0206 0.0437 0.0012 0.0488
AG 0.13 0.41 1.84 0.0021 0.0014 0.151 0.0018 0.2737 0.3965 0.3658 0.0590 0.0019 0.5168 0.0010 0.0394
AH 0.18 0.88 1.99 0.0104 0.0011 0.050 0.0124 0.0290 0.0309 0.0557 0.1003 0.0010 0.0287 0.0102 0.3065
AI 0.12 1.07 1.58 0.0011 0.0023 0.089 0.0015 0.0574 0.0245 0.0368 0.0828 0.0025 0.0481 0.0012 0.5162
AJ 0.17 0.34 1.90 0.0056 0.0108 0.735 0.0023 0.0361 0.0414 0.0298 0.0347 0.0012 0.0498 0.0008 0.4162
AK 0.15 1.02 2.17 0.0012 0.0013 0.065 0.0019 0.0346 0.0265 0.3168 0.0353 0.0164 0.2737 0.0055 0.0373
AL 0.09 0.77 1.41 0.0092 0.0014 0.189 0.0068 0.0292 0.4040 0.0816 0.0900 0.0008 0.0323 0.0013 0.0411
AM 0.16 0.15 2.27 0.0012 0.0057 0.177 0.0017 0.1965 0.0378 0.0683 0.1296 0.0013 0.0255 0.0010 0.0529
AN 0.19 0.52 2.72 0.0023 0.0146 0.104 0.0147 0.0553 0.0970 0.0589 0.1131 0.0020 0.0213 0.0025 0.0646
AO 0.18 0.43 1.86 0.0015 0.0018 0.176 0.0020 0.0352 0.3635 0.1409 0.7894 0.0026 0.2141 0.0009 0.1646
AP 0.14 1.18 2.38 0.0014 0.0065 0.248 0.0009 0.1326 0.0298 0.0320 0.0450 0.0017 0.0308 0.0012 0.0618
AQ 0.07 0.49 1.39 0.0009 0.0012 0.529 0.0054 0.0448 0.2735 0.0714 0.0543 0.0011 0.0256 0.0072 0.0265
AR 0.12 0.28 2.02 0.0010 0.0023 0.116 0.0019 0.0356 0.0886 0.0769 0.0631 0.0022 0.0275 0.0079 0.2734
AS 0.06 0.55 2.66 0.0031 0.0034 0.092 0.0010 0.0661 0.0500 0.0342 0.0752 0.0014 0.0449 0.0006 0.0384
AT 0.06 0.53 1.25 0.0147 0.0048 0.146 0.0013 0.0304 0.3045 0.2300 0.1038 0.0014 0.0567 0.0031 0.0300
AU 0.07 1.11 2.04 0.0114 0.0020 0.084 0.0017 0.4196 0.0236 0.3852 0.1078 0.0014 0.0190 0.0007 0.0392
AV 0.20 1.01 2.53 0.0011 0.0017 0.155 0.0076 0.0588 0.0558 0.4262 0.1542 0.0027 0.0448 0.0005 0.1713
AW 0.11 0.65 2.49 0.0081 0.0022 0.028 0.0034 - - 0.0503 0.5112 0.0012 0.0213 0.0021 0.0108
Bold and underlined values are outside the scope of the present invention.
[0097]
[Table 1-3]
35
Table 1-3
No. V Cu W Ta Sn Sb As Mg Ca Y Zr La Ce Note
A 0.1842 0.0197 0.0116 0.0101 0.0039 0.0054 0.0048 0.0054 0.0026 0.0054 0.0116 0.0208 0.0421 developed steel
B 0.0187 0.0500 0.0052 0.0080 0.0393 0.0069 0.0315 0.0042 0.0409 0.0046 0.0051 0.0032 0.0027 developed steel
C 0.1030 0.1234 0.0074 0.0080 0.0046 0.0198 0.0053 0.0057 0.0117 0.0041 0.0307 0.0027 0.0043 developed steel
D 0.0304 0.0260 0.0802 0.0063 0.0034 0.0304 0.0086 0.0050 0.0021 0.0050 0.0043 0.0397 0.0394 developed steel
E - - - - - - - - - - - - - developed steel
F 0.0376 0.0377 0.0056 0.0093 0.0041 0.0046 0.0382 0.0038 0.0189 0.0063 0.0193 0.0026 0.0031 developed steel
G - - - - - - - - - - - - - developed steel
H 0.0486 0.4274 0.0097 0.0825 0.0068 0.0388 0.0030 0.0057 0.0056 0.0018 0.0037 0.0110 0.0033 developed steel
I 0.4001 0.0561 0.0858 0.0786 0.0414 0.0038 0.0026 0.0069 0.0401 0.0081 0.0072 0.0052 0.0041 developed steel
J 0.0687 0.3056 0.0041 0.0089 0.0039 0.0043 0.0046 0.0040 0.0039 0.0108 0.0028 0.0055 0.0044 developed steel
K - - - - - - - - - - - - - developed steel
L - - - - - - - - - - - - - developed steel
M 0.0220 0.0384 0.0086 0.0095 0.0070 0.0040 0.0045 0.0133 0.0097 0.0199 0.0043 0.0032 0.0055 developed steel
N - - - - - - - - - - - - - developed steel
O 0.0630 0.0251 0.0103 0.0054 0.0049 0.0044 0.0045 0.0307 0.0050 0.0021 0.0058 0.0025 0.0023 developed steel
P 0.3300 0.0430 0.0059 0.0113 0.0311 0.0082 0.0067 0.0080 0.0043 0.0428 0.0056 0.0068 0.0118 developed steel
Q - - - - - - - - - - - - - developed steel
R 0.0333 0.3815 0.0227 0.0121 0.0208 0.0019 0.0049 0.0429 0.0020 0.0031 0.0027 0.0039 0.0181 developed steel
S - - - - - - - - - - - - - comparative steel
T - - - - - - - - - - - - - comparative steel
U - - - - - - - - - - - - - comparative steel
V - - - - - - - - - - - - - comparative steel
W - - - - - - - - - - - - - comparative steel
X - - - - - - - - - - - - - comparative steel
Bold and underlined values are outside the scope of the present invention.
[0098]
5 [Table 1-4]
P190781WO
36
Table 1-4
No. V Cu W Ta Sn Sb As Mg Ca Y Zr La Ce Note
Y 0.0607 0.2154 0.0497 0.0810 0.0043 0.0041 0.0363 0.0033 0.0042 0.0048 0.0329 0.0082 0.0080 comparative steel
Z - - - - - - - - - - - - - comparative steel
AA 0.0497 0.0377 0.0060 0.0284 0.0104 0.0060 0.0041 0.0063 0.0082 0.0404 0.0109 0.0375 0.0256 comparative steel
AB 0.4086 0.1029 0.0068 0.0738 0.0059 0.0075 0.0419 0.0035 0.0368 0.0042 0.0056 0.0091 0.0030 comparative steel
AC 0.0368 0.0320 0.0132 0.0064 0.0027 0.0058 0.0033 0.0365 0.0231 0.0392 0.0106 0.0025 0.0091 comparative steel
AD 0.0580 0.0529 0.0070 0.0088 0.0362 0.0377 0.0037 0.0043 0.0042 0.0224 0.0067 0.0071 0.0043 comparative steel
AE 0.0443 0.0418 0.0072 0.0057 0.0036 0.0032 0.0334 0.0062 0.0050 0.0063 0.0100 0.0114 0.0050 comparative steel
AF 0.0437 0.1359 0.0061 0.0106 0.0046 0.0080 0.0250 0.0403 0.0025 0.0052 0.0051 0.0077 0.0046 comparative steel
AG 0.1018 0.0379 0.0070 0.0090 0.0269 0.0068 0.0037 0.0046 0.0155 0.0050 0.0414 0.0043 0.0072 comparative steel
AH 0.0353 0.0704 0.0161 0.0068 0.0039 0.0032 0.0306 0.0025 0.0425 0.0061 0.0166 0.0032 0.0139 comparative steel
AI 0.0953 0.0520 0.0232 0.0155 0.0041 0.0025 0.0080 0.0028 0.0052 0.0183 0.0044 0.0142 0.0029 comparative steel
AJ 0.5123 0.0518 0.0105 0.0060 0.0125 0.0361 0.0142 0.0028 0.0052 0.0040 0.0050 0.0031 0.0047 comparative steel
AK 0.0758 0.5128 0.0428 0.0038 0.0304 0.0052 0.0022 0.0146 0.0030 0.0043 0.0023 0.0132 0.0428 comparative steel
AL 0.0249 0.0480 0.1022 0.0802 0.0406 0.0049 0.0050 0.0047 0.0032 0.0040 0.0218 0.0028 0.0023 comparative steel
AM 0.0295 0.0524 0.0731 0.1034 0.0041 0.0260 0.0041 0.0216 0.0104 0.0052 0.0026 0.0044 0.0037 comparative steel
AN 0.0383 0.0191 0.0257 0.0094 0.0516 0.0170 0.0033 0.0070 0.0377 0.0032 0.0039 0.0026 0.0069 comparative steel
AO 0.2713 0.3775 0.0084 0.0049 0.0025 0.0519 0.0086 0.0034 0.0050 0.0345 0.0391 0.0035 0.0034 comparative steel
AP 0.0358 0.0613 0.0034 0.0117 0.0048 0.0392 0.0515 0.0037 0.0028 0.0025 0.0194 0.0210 0.0061 comparative steel
AQ 0.4039 0.2615 0.0053 0.0074 0.0060 0.0123 0.0190 0.0510 0.0401 0.0044 0.0040 0.0065 0.0026 comparative steel
AR 0.0405 0.0822 0.0054 0.0060 0.0054 0.0432 0.0396 0.0037 0.0518 0.0108 0.0053 0.0028 0.0044 comparative steel
AS 0.1617 0.0326 0.0084 0.0042 0.0365 0.0066 0.0018 0.0039 0.0051 0.0517 0.0382 0.0030 0.0042 comparative steel
AT 0.0368 0.0332 0.0380 0.0452 0.0076 0.0024 0.0041 0.0049 0.0114 0.0039 0.0513 0.0024 0.0043 comparative steel
AU 0.0430 0.0742 0.0111 0.0103 0.0056 0.0023 0.0057 0.0103 0.0038 0.0118 0.0028 0.0519 0.0057 comparative steel
AV 0.3857 0.0309 0.0105 0.0115 0.0048 0.0023 0.0026 0.0026 0.0041 0.0027 0.0047 0.0052 0.0511 comparative steel
AW - - - - - - - - - - - - - developed steel
Bold and underlined values are outside the scope of the present invention.
[0099]
5 [Table 2-1]
P190781WO
37
Table 2-1
No.
Total of
ferrite and
bainite
(%)
Total of martensite and
tempered martensite
(%)
Total of pearlite
and retained
austenite
(%)
Integration degree of
(111)<112> orientation
of ferrite
Integration degree of
(252)<2-11> orientation of
martensite and tempered
martensite
Tensile
strength
(MPa)
Total
elongation
(%)
Stretch formability Note
A-1 41.9 57.1 1.0 3.8 0.6 996 15.2 ○ Example
B-1 43.9 55.0 1.1 5.2 0.3 1,098 13.5 ○ Example
C-1 50.8 48.0 1.2 4.8 2.4 1,045 14.6 ○ Example
D-1 53.8 44.3 1.9 5.4 2.9 1,030 14.8 ○ Example
E-1 89.2 6.0 4.8 3.9 1.2 594 29.4 ○ Example
F-1 83.3 6.3 10.4 5.9 1.5 623 27.3 ○ Example
G-1 50.2 48.2 1.6 5.1 0.4 988 15.5 ○ Example
H-1 49.5 49.0 1.5 5.9 0.5 1,071 14.2 ○ Example
I-1 68.2 28.2 3.6 4.3 2.6 903 16.8 ○ Example
J-1 40.0 58.3 1.7 4.8 4.8 1,023 14.9 ○ Example
K-1 76.3 18.9 4.8 4.7 2.2 899 17.2 ○ Example
L-1 39.3 59.6 1.1 4.9 1.1 1,079 14.2 ○ Example
M-1 87.2 10.5 2.3 5.6 4.5 895 17.1 ○ Example
N-1 70.3 27.0 2.7 3.3 0.3 786 19.7 ○ Example
O-1 76.0 20.8 3.2 3.2 2.1 1,011 14.9 ○ Example
P-1 75.0 24.0 1.0 5.4 1.1 1,032 14.7 ○ Example
Q-1 75.9 21.6 2.5 3.5 4.3 991 15.0 ○ Example
R-1 48.0 51.0 1.0 4.3 1.2 1,035 14.6 ○ Example
S-1 89.2 7.2 3.6 3.7 5.6 632 24.1  Comparative Example
T-1 22.6 76.4 1.0 2.4 2.2 896 14.1  Comparative Example
U-1 42.9 56.1 1.0 3.5 0.6 1,164 11.8  Comparative Example
V-1 80.2 4.1 15.7 5.5 4.9 488 12.7 ○ Comparative Example
W-1 7.2 79.6 13.2 5.0 2.2 1,077 13.5  Comparative Example
X-1 82.1 10.1 7.8 5.1 4.8 723 21.9  Comparative Example
Bold and underlined values are outside the scope of the present invention.
[0100]
5 [Table 2-2]
P190781WO
38
Table 2-2
No.
Total of
ferrite and
bainite
(%)
Total of martensite and
tempered martensite
(%)
Total of pearlite
and retained
austenite
(%)
Integration degree of
(111)<112> orientation
of ferrite
Integration degree of
(252)<2-11> orientation of
martensite and tempered
martensite
Tensile
strength
(MPa)
Total
elongation
(%)
Stretch formability Note
Y-1 68.9 28.4 2.7 5.5 2.1 848 18.1  Comparative Example
Z-1 91.6 7.3 1.1 5.2 3.7 538 21.2 ○ Comparative Example
AA-1 81.8 14.2 4.0 3.2 4.1 800 19.4  Comparative Example
AB-1 65.1 33.1 1.8 3.2 1.1 1,001 15.1  Comparative Example
AC-1 58.0 40.5 1.5 4.2 3.8 1,098 13.8  Comparative Example
AD-1 11.2 87.8 1.0 5.3 2.3 1,116 10.4  Comparative Example
AE-1 10.8 72.9 16.3 4.3 4.9 980 19.3  Comparative Example
AF-1 59.1 39.8 1.1 5.1 0.3 1,000 13.2  Comparative Example
AG-1 50.4 48.1 1.5 3.6 3.2 1,007 12.3  Comparative Example
AH-1 55.9 42.1 2.0 5.6 0.2 1,095 11.4  Comparative Example
AI-1 75.2 20.9 3.9 5.7 3.7 988 14.6  Comparative Example
AJ-1 69.5 28.5 2.0 4.9 4.2 981 13.8  Comparative Example
AK-1 47.7 51.2 1.1 3.2 0.3 1,117 12.8  Comparative Example
AL-1 80.9 14.8 4.3 4.5 3.4 842 16.9  Comparative Example
AM-1 55.8 41.9 2.3 3.6 1.1 957 13.4  Comparative Example
AN-1 37.4 61.4 1.2 3.7 4.3 1,034 8.9  Comparative Example
AO-1 43.8 55.0 1.2 3.1 1.1 1,040 9.2  Comparative Example
AP-1 59.2 39.6 1.2 5.2 1.3 1,089 8.1  Comparative Example
AQ-1 87.8 7.0 5.2 4.4 1.5 731 11.4  Comparative Example
AR-1 68.8 27.7 3.5 4.1 2.5 867 10.6  Comparative Example
AS-1 64.4 34.2 1.4 3.7 3.6 921 12.2  Comparative Example
AT-1 85.9 9.9 4.2 4.0 3.7 717 11.5  Comparative Example
AU-1 64.1 34.8 1.1 5.8 1.4 1,032 6.3  Comparative Example
AV-1 40.1 58.9 1.0 5.4 4.7 1,106 7.8  Comparative Example
AW-1 23.4 76.6 0.0 5.2 3.8 1,078 13.1 ○ Example
Bold and underlined values are outside the scope of the present invention.
P190781WO
39
[0101]
Referring to Table 2, in Example S-1, the orientation integration degree of martensite was
not randomized due to the low C content, and the integration degree of the (252)<2-11>
orientation of martensite and tempered martensite was thus higher than 5.0. As a result, the
stretch formability was deteriorated. In Example 5 T-1, since the orientation integration degree of
ferrite was reduced due to the high C content, the stretch formability was deteriorated. In
Example U-1, due to the high Si content, the tensile strength was increased to cause
embrittlement, and the stretch formability was deteriorated. In Example V-1, the tensile strength
was reduced due to the low Mn content. In Example W-1, since ferrite transformation and bainite
10 transformation were inhibited due to the high Mn content, the stretch formability was
deteriorated. In Example X-1, due to the high P content, the steel sheet was embrittled and the
stretch formability was deteriorated. In Example Y-1, due to the high S content, cracking
occurred during cold forming, and the stretch formability was deteriorated. In Example Z-1,
since ferrite transformation and bainite transformation were excessively facilitated due to the
15 high Al content, the tensile strength was reduced. In Example AA-1, due to the high N content,
coarse nitrides were formed in the steel sheet, and the stretch formability was deteriorated.
[0102]
In Example AB-1, a large amount of fine Co carbides precipitated due to the high Co
content, and the stretch formability was deteriorated. In Example AC-1, the stretch formability
20 was deteriorated due to the high Ni content. In Example AD-1, since martensitic transformation
was facilitated due to the high Mo content, the stretch formability was deteriorated. In Example
AE-1, a large amount of retained austenite was generated due to the high Cr content, and the
stretch formability was deteriorated. In Example AF-1, oxides were formed due to the high O
content, and the stretch formability was deteriorated. In Example AG-1, precipitation of
25 carbonitride was increased due to the high Ti content, and the stretch formability was
deteriorated. In Example AH-1, since coarse B oxides were generated in the steel due to the high
B content, the stretch formability was deteriorated. In Example AI-1, Nb carbide precipitated in a
large amount due to the high Nb content, and the stretch formability was deteriorated. In
Example AJ-1, a large amount of carbonitride precipitated due to the high V content, and the
30 stretch formability was deteriorated.
[0103]
In Example AK-1, the tensile strength was excessively high due to the high Cu content, and
the stretch formability was deteriorated in association therewith. In Example AL-1, the stretch
formability was deteriorated due to the high W content. In Example AM-1, a large amount of
35 fine Ta carbide precipitated due to the high Ta content, and the stretch formability was
deteriorated. In Example AN-1, embrittlement of ferrite occurred due to the high Sn content, and
40
the stretch formability was thereby deteriorated. In Examples AO-1 and AP-1, grain boundary
segregation occurred due to the high Sb content and the high As content, respectively, and the
stretch formability was thereby deteriorated. In Example AQ-1, coarse inclusions were formed
due to the high Mg content, and the stretch formability was thereby deteriorated. In Example
AR-1, the stretch formability was deteriorated du 5 e to the high Ca content. In Examples AS-1 to
AV-1, coarse oxides were generated due to the high content of Y, Zr, La and Ce, respectively,
and the stretch formability was deteriorated.
[0104]
In contrast to the above, in Examples A-1 to R-1, steel sheets having a high strength and
10 excellent stretch formability were obtained by appropriately controlling the chemical
composition and the structure of each steel sheet as well as the integration degrees of ferrite and
martensite.
[0105]
[Example 2]
15 Further, in order to investigate the effects of the production conditions, hot-rolled steel
sheets of 2.3 mm in thickness were produced by performing thermo-mechanical treatments in
accordance with the production conditions shown in Table 3 on the respective steel species A to
R that had been confirmed to have excellent properties as shown in Table 2, and the properties of
these steel sheets were evaluated after cold rolling and annealing. The symbols GI and GA under
20 “Plating treatment” each indicate a method of galvanization treatment. The symbol GI represents
a steel sheet which was immersed in a 460C hot-dip galvanizing bath and thereby provided with
a galvanized layer on the surface, and the symbol GA represents a steel sheet which was
immersed in a hot-dip galvanizing bath, subsequently heated to 485C, and thereby provided
with an alloy layer of iron and zinc on the surface. In addition, on each of the steel sheets, a
25 tempering treatment, in which the steel sheet once cooled to 150C was reheated and retained for
2 to 120 seconds in a period between retention of the steel sheet at the respective retention
temperatures in cold-rolled sheet annealing and subsequent cooling of the steel sheet to room
temperature, was performed. It is noted here that, in those Examples where the tempering time
was in a range of 3,600 to 33,000 seconds, each steel sheet wound into a coil form after being
30 cooled to room temperature was tempered using a separate annealing apparatus (box annealing
furnace). Moreover, in those Examples with a description of “none” for tempering in Table 3,
tempering was not performed. The thus obtained results are shown in Table 4. As for the
methods of evaluating the properties, the same methods were employed as in Example 1.
[0106]
35 [Table 3-1]
P190781WO
41
Table 3-1
No.
Slab
rolling
reduction
ratio
(%)
Hot-rolling
finishing
temperature
(C)
Coiling
temperature
(C)
Retention
time at
[coiling
start
temperature
+ 20C to
100C]
(min)
Cold-rolling
reduction
ratio
(%)
Cold-rolled sheet annealing
Plating
treatment
Cold-rolled sheet annealing
Skin
pass
rolling
ratio
(%)
Note
Heating
rate
(C/sec)
Annealing
temperature
(C)
Retention
time
(sec)
Average
cooling
rate
(C/sec)
Cooling
stop
temperature
(C)
Retention
temperature
(C)
Retention
time
(sec)
Tempering
temperature
(C)
Tempering
time
(sec)
A-2 11 910 432 134 72.5 58.2 821 83 63.6 429 362 314 none 212 20 0.3 Example
B-2 18 796 484 204 26.5 41.2 757 547 10.1 504 504 287 GA 334 19 0.5 Example
C-2 11 685 416 134 36.0 46.3 808 55 69.6 514 399 65 none 343 47 0.3 Example
D-2 21 726 677 174 93.2 21.1 746 564 62.4 484 426 146 none 180 24 0.3 Comparative Example
E-2 10 864 599 154 6.8 39.7 898 386 83.5 316 535 450 none 325 34 0.3 Comparative Example
F-2 46 877 680 41 76.8 72.4 853 307 97.7 286 378 437 none 350 13 0.5 Comparative Example
G-2 35 850 657 103 30.2 34.8 738 331 8.4 187 382 61 none 208 3,600 0.6 Example
H-2 15 740 508 103 86.8 85.2 800 385 41.9 164 448 217 none 223 35 0.4 Example
I-2 38 728 566 284 47.3 69.8 743 117 27.4 386 511 238 none 315 62 0.2 Example
J-2 36 686 549 169 62.4 25.2 728 118 11.7 530 529 372 GA 202 27 0.2 Example
K-2 31 696 479 86 28.1 50.6 708 576 35.4 532 472 554 none 175 11 0.6 Example
L-2 21 851 677 3 31.9 54.7 802 173 22.1 132 386 195 none 214 43 0.4 Comparative Example
M-2 28 753 635 72 64.2 25.5 844 408 93.4 312 467 206 none 311 27 0.1 Example
N-2 32 933 494 68 89.1 82.9 820 528 76.0 531 528 117 GA 375 252 0.1 Example
O-2 26 785 447 139 18.8 90.3 775 309 24.3 83 525 337 none 256 7 0.6 Example
P-2 20 827 634 220 23.4 5.7 724 434 63.8 94 424 404 none 234 20 0.3 Example
Q-2 16 767 590 276 42.6 97.0 919 66 79.8 191 413 320 none 196 55 0.2 Comparative Example
R-2 26 637 429 28 15.9 15.2 737 213 38.3 285 492 512 none 162 3 0.2 Comparative Example
A-3 9 897 618 252 23.1 30.1 763 141 82.6 235 419 359 none none none 0.4 Example
B-3 28 794 585 68 69.0 52.8 745 184 18.9 284 436 396 none 232 12 0.1 Example
C-3 13 696 611 182 28.5 15.7 846 267 8.9 406 537 100 none 186 24 0.1 Example
D-3 15 829 475 257 65.2 88.0 820 515 66.8 512 512 332 GI 377 23 0.5 Example
E-3 7 817 597 70 88.2 11.2 828 126 80.1 445 525 211 none 368 29 0.5 Example
F-3 35 913 525 29 72.5 62.7 870 373 90.6 405 387 334 none 191 5 0.4 Example
G-3 28 667 658 119 81.3 61.0 716 422 46.3 545 533 419 GI 263 119 0.3 Example
H-3 25 690 668 225 43.5 79.6 808 484 19.4 237 355 137 none 272 33 0.5 Example
I-3 37 745 554 245 47.1 59.8 807 445 46.9 521 501 263 GA 340 24 0.4 Example
Bold and underlined values are outside the scope of the present invention.
[0107]
5 [Table 3-2]
P190781WO
42
Table 3-2
No.
Slab
rolling
reduction
ratio
(%)
Hot-rolling
finishing
temperature
(C)
Coiling
temperature
(C)
Retention
time at
[coiling start
temperature +
20C to
100C]
(min)
Cold-rolling
reduction
ratio
(%)
Cold-rolled sheet annealing
Plating
treatment
Cold-rolled sheet annealing
Skin
pass
rolling
ratio
(%)
Note
Heating
rate
(C/sec)
Annealing
temperature
(C)
Retention
time
(sec)
Average
cooling
rate
(C/sec)
Cooling
stop
temperature
(C)
Retention
temperature
(C)
Retention
time
(sec)
Tempering
temperature
(C)
Tempering
time
(sec)
J-3 14 782 503 15 53.9 67.3 796 27 49.5 352 518 59 none 341 36 0.3 Example
K-3 11 673 426 217 71.0 87.1 721 341 53.1 210 368 123 none 278 33,000 0.4 Example
L-3 39 883 446 269 44.1 38.1 737 297 91.9 362 425 288 none 209 14 0.4 Example
M-3 25 826 405 204 85.5 25.8 866 464 74.2 500 358 152 none 336 11 0.6 Example
N-3 36 904 466 244 51.2 13.9 845 79 28.8 547 547 403 GA 297 13 0.3 Example
O-3 31 713 503 174 20.6 6.1 835 254 68.4 158 494 201 none 258 22 0.5 Example
P-3 23 969 568 48 15.3 64.2 828 560 76.1 534 543 303 none 245 55 0.2 Comparative Example
Q-3 30 886 580 280 75.2 94.9 719 59 26.5 47 440 446 none 356 28 0.4 Example
R-3 19 754 733 187 58.5 65.5 877 505 76.5 233 528 30 none 266 15 0.4 Comparative Example
A-4 20 840 641 35 76.8 80.4 790 505 53.7 509 496 399 GI none none 0.2 Example
B-4 6 794 673 185 37.7 19.0 777 246 35.1 479 476 593 none 308 43 0.5 Example
C-4 23 931 448 317 35.2 48.7 893 120 59.9 468 413 513 none 276 58 0.2 Comparative Example
D-4 29 782 524 223 21.0 76.1 760 338 8.4 140 478 165 none 160 49 0.4 Example
E-4 7 710 391 19 59.7 48.6 857 365 50.7 205 490 531 none 246 6 0.3 Comparative Example
F-4 12 716 565 84 41.0 6.7 776 226 47.2 400 450 163 none 347 381 0.3 Example
G-4 38 720 448 280 46.3 56.6 841 148 58.7 126 445 490 none 303 55 0.3 Example
H-4 17 865 552 115 38.3 8.5 819 37 69.3 170 399 453 none 388 19,000 0.2 Example
I-4 9 939 455 120 60.3 45.7 692 152 88.0 487 437 235 none 367 31 0.5 Comparative Example
J-4 21 766 644 152 35.2 29.9 882 334 54.9 149 499 247 none 293 10 0.2 Example
K-4 32 899 537 59 69.2 75.7 750 540 16.3 315 373 556 none 302 23 0.3 Example
L-4 18 814 630 113 54.3 9.1 756 484 30.5 398 399 484 none 287 17 0.5 Example
M-4 33 848 632 190 13.7 72.6 777 215 39.6 512 492 532 GA 298 61 0.5 Example
N-4 14 931 615 52 18.9 40.3 791 252 24.0 360 516 81 none 375 39 0.2 Example
O-4 4 768 526 211 79.3 78.8 872 358 84.4 467 444 47 none 382 57 0.4 Comparative Example
P-4 23 713 513 264 60.4 36.2 783 437 15.7 541 537 171 GA none none 0.2 Example
Q-4 8 878 465 84 54.5 22.7 722 290 37.6 423 460 559 none 178 39 0.4 Example
R-4 28 669 545 280 36.9 84.3 878 181 86.7 223 426 476 none 164 43 0.5 Example
Bold and underlined values are outside the scope of the present invention.
[0108]
5 [Table 4-1]
P190781WO
43
Table 4-1
No.
Total of ferrite and
bainite
(%)
Total of martensite
and tempered
martensite
(%)
Total of
pearlite and
retained
austenite
(%)
Integration degree of
(111)<112>
orientation of ferrite
Integration degree of
(252)<2-11> orientation of
martensite and tempered
martensite
Tensile
strength
(MPa)
Total
elongation
(%)
Stretch
formability
Sheet
thickness
(mm)
Note
A-2 19.3 79.6 1.1 3.9 1.2 891 17.3 ○ 0.6 Example
B-2 57.9 40.8 1.3 4.7 2.2 1,047 14.6 ○ 1.7 Example
C-2 28.7 70.3 1.0 5.1 4.0 963 15.9 ○ 1.5 Example
D-2 56.5 42.8 0.7 4.5 5.4 1,037 14.6  0.2 Comparative Example
E-2 12.6 85.6 1.8 2.3 2.7 1,040 14.3  2.1 Comparative Example
F-2 23.7 75.2 1.1 2.8 1.3 891 17.3  0.5 Comparative Example
G-2 65.4 31.3 3.3 4.0 1.8 777 20.2 ○ 1.6 Example
H-2 31.7 67.3 1.0 3.8 3.3 984 15.5 ○ 0.3 Example
I-2 69.3 29.1 1.6 4.4 4.7 927 16.5 ○ 1.2 Example
J-2 59.4 39.2 1.4 5.1 1.6 1,036 14.8 ○ 0.9 Example
K-2 77.2 21.8 1.0 3.9 1.2 808 19.2 ○ 1.7 Example
L-2 26.3 73.3 0.4 6.1 5.8 1,058 14.1  1.6 Comparative Example
M-2 58.3 40.7 1.0 6.0 0.6 981 15.7 ○ 0.8 Example
N-2 49.1 49.8 1.1 3.4 0.4 840 18.1 ○ 0.3 Example
O-2 73.8 25.0 1.2 3.8 4.2 1,044 14.4 ○ 1.9 Example
P-2 87.3 11.7 1.0 3.2 4.9 1,036 14.5 ○ 1.8 Example
Q-2 11.9 78.3 9.8 2.4 1.6 984 15.3  1.3 Comparative Example
R-2 71.1 27.9 1.0 1.1 6.8 1,001 14.9  1.9 Comparative Example
A-3 47.4 51.4 1.2 4.5 1.9 992 15.4 ○ 1.8 Example
B-3 64.4 34.5 1.1 5.0 1.0 1,052 14.3 ○ 0.7 Example
C-3 19.0 78.9 2.1 3.3 1.9 882 17.3 ○ 1.6 Example
D-3 13.9 79.8 6.3 4.0 1.6 930 16.4 ○ 0.8 Example
E-3 32.7 65.6 1.7 3.6 3.9 1,100 13.4 ○ 0.3 Example
F-3 23.6 75.2 1.2 5.3 3.0 891 16.9 ○ 0.6 Example
G-3 73.6 25.0 1.4 4.3 4.7 814 19.1 ○ 0.4 Example
H-3 30.1 68.9 1.0 4.5 2.1 971 15.8 ○ 1.3 Example
I-3 36.5 62.4 1.1 5.2 2.1 974 15.8 ○ 1.2 Example
Bold and underlined values are outside the scope of the present invention.
[0109]
5 [Table 4-2]
P190781WO
44
Table 4-2
No.
Total of
ferrite and
bainite
(%)
Total of martensite
and tempered
martensite
(%)
Total of
pearlite and
retained
austenite
(%)
Integration degree of
(111)<112>
orientation of ferrite
Integration degree of
(252)<2-11> orientation
of martensite and
tempered martensite
Tensile
strength
(MPa)
Total
elongation
(%)
Stretch
formability
Sheet
thickness
(mm)
Note
J-3 15.3 79.6 5.1 5.4 3.0 902 16.6 ○ 1.1 Example
K-3 83.9 12.7 3.4 5.0 4.5 761 20.3 ○ 0.7 Example
L-3 60.8 38.2 1.0 4.2 1.3 1,031 14.8 ○ 1.3 Example
M-3 52.4 46.6 1.0 5.4 1.6 991 15.2 ○ 0.3 Example
N-3 47.9 48.7 3.4 5.9 3.2 760 20.4 ○ 1.1 Example
O-3 44.7 54.3 1.0 4.5 0.9 1,064 14 ○ 1.8 Example
P-3 58.1 40.9 1.0 1.5 6.1 1,060 14.3  1.9 Comparative Example
Q-3 89.1 9.2 1.7 4.5 1.5 945 15.8 ○ 0.6 Example
R-3 13.8 78.9 7.3 5.5 7.5 923 16.5  1 Comparative Example
A-4 34.5 64.5 1.0 4.7 3.8 984 15.3 ○ 0.5 Example
B-4 45.5 53.5 1.0 5.7 4.6 1,024 15 ○ 1.4 Example
C-4 - - - - - - - - - -
D-4 55.3 43.6 1.1 4.3 2.1 1,036 14.5 ○ 1.8 Example
E-4 26.3 72.6 1.1 3.2 6.7 1,085 13.9  0.9 Comparative Example
F-4 61.3 37.5 1.2 6.1 3.9 827 18.7 ○ 1.4 Example
G-4 16.7 78.7 4.6 3.5 0.4 924 16.2 ○ 1.2 Example
H-4 19.6 79.4 1.0 5.9 1.0 860 18 ○ 1.4 Example
I-4 88.6 0.3 11.1 6.1 3.8 468 24.5  0.9 Comparative Example
J-4 16.4 79.2 4.4 4.4 1.1 911 16.9 ○ 1.5 Example
K-4 77.6 21.4 1.0 6.1 1.3 833 18.6 ○ 0.7 Example
L-4 50.5 48.4 1.1 3.8 2.5 1,018 14.6 ○ 1.1 Example
M-4 81.4 17.6 1.0 3.3 0.8 928 16.2 ○ 2 Example
N-4 64.5 34.4 1.1 4.9 2.5 717 22 ○ 1.9 Example
O-4 27.0 72.0 1.0 2.2 5.9 988 15.3  0.5 Comparative Example
P-4 74.1 24.9 1.0 5.1 4.3 1,033 14.6 ○ 0.9 Example
Q-4 89.4 9.2 1.4 4.8 0.4 948 15.8 ○ 1 Example
R-4 14.4 79.8 5.8 3.9 1.5 911 16.5 ○ 1.5 Example
Bold and underlined values are outside the scope of the present invention.
P190781WO
45
[0110]
Referring to Table 4, in Example D-2, the integration degree of the (252)<2-11> orientation
of martensite and tempered martensite was high due to the high rolling reduction ratio in the cold
rolling, as a result of which the stretch formability was deteriorated. In Example E-2, the
integration degree of the (111)<5 112> orientation was low due to the low rolling reduction ratio
in the cold rolling, as a result of which the stretch formability was deteriorated. In Example F-2,
since the rolling reduction ratio in the casting step was excessively high, the integration degree
of the (111)<112> orientation of ferrite after the cold-rolled sheet annealing was low, and the
stretch formability was consequently deteriorated. In Example L-2, since the post-coiling
10 retention time at the prescribed temperature was short, the integration degree of the (252)<2-11>
orientation of martensite and tempered martensite was not reduced, as a result of which the
stretch formability was deteriorated.
[0111]
In Example Q-2, since the annealing temperature was high, the integration degree of the
15 (111)<112> orientation of ferrite was low, as a result of which the stretch formability was
deteriorated. In Example R-2, since the hot-rolling finishing temperature was low, a rolled
crystallographic texture of austenite developed, and this caused anisotropy in the steel material
properties, as a result of which the integration degree of the (252)<2-11> orientation of
martensite in the final product was not reduced, and the stretch formability was deteriorated. In
20 Example P-3, since the hot-rolling finishing temperature was high, abnormal growth of austenite
grains occurred and the crystallographic texture could not be made isotropic, as a result of which
the integration degree of the (111)<112> orientation of ferrite was reduced, and the stretch
formability was deteriorated. In Example R-3, due to the high coiling temperature, pearlite
transformation was allowed to progress in the post-coiling heating treatment, and the target
25 hot-rolled structure was thus not obtained, as a result of which the integration degree of the
(252)<2-11> orientation of martensite in the final product was high, and the stretch formability
was deteriorated.
[0112]
In Example C-4, since the post-coiling retention time at the prescribed temperature was
30 long, internal oxides were formed in the hot-rolled sheet, and cracks were generated on the steel
sheet surface in the subsequent treatments. Therefore, the analysis of the structure and the
evaluation of the mechanical properties were not performed. In Example E-4, due to the low
coiling temperature, the target hot-rolled structure was not obtained even with a post-coiling
heating treatment, as a result of which the integration degree of the (252)<2-11> orientation of
35 martensite in the final product was high, and the stretch formability was deteriorated. In Example
I-4, due to the low annealing temperature, not only the amount of generated austenite was small
46
and the ratio of martensite structure was reduced in the structure after the cold rolling and
annealing, but also non-recrystallized ferrite remained, as a result of which the tensile strength
and the stretch formability were both deteriorated. In Example O-4, due to the low rolling
reduction ratio in the casting step, the integration degree of the (111)<112> orientation of ferrite
was low and the integration de 5 gree of the (252)<2-11> orientation of martensite and tempered
martensite was high, as a result of which the stretch formability was deteriorated.
[0113]
In contrast to the above, in all of Examples according to the present invention, a steel sheet
having a high strength and excellent stretch formability was obtained particularly by performing
10 rolling reduction at a prescribed reduction ratio in the casting step, and additionally controlling
the hot-rolling finishing temperature, the coiling, the cold rolling, and the annealing as
appropriate.
[0114]
FIG. 1 is a graph showing the effects of the integration degree of the (111)<112>
15 orientation of ferrite and the integration degree of the (252)<2-11> orientation of martensite and
tempered martensite on the stretch formability of the DP steels used in Examples 1 and 2. As
apparent from FIG. 1, it is understood that a steel sheet having excellent stretch formability can
be obtained by controlling the integration degree of the (111)<112> orientation of ferrite to be
3.0 or higher, and the integration degree of the (252)<2-11> orientation of martensite and
20 tempered martensite to be 5.0 or lower.

WE CLAIMS

A steel sheet, having a chemical composition comprising, by mass %:
5 C: 0.05 to 0.20%;
Si: 0.01 to 1.30%;
Mn: 1.00 to 3.00%;
P: 0.0001 to 0.0200%;
S: 0.0001 to 0.0200%;
10 Al: 0.001 to 1.000%;
N: 0.0001 to 0.0200%;
Co: 0 to 0.5000%;
Ni: 0 to 0.5000%;
Mo: 0 to 0.5000%;
15 Cr: 0 to 1.0000%;
O: 0 to 0.0200%;
Ti: 0 to 0.5000%;
B: 0 to 0.0100%;
Nb: 0 to 0.5000%;
20 V: 0 to 0.5000%;
Cu: 0 to 0.5000%;
W: 0 to 0.1000%;
Ta: 0 to 0.1000%;
Sn: 0 to 0.0500%;
25 Sb: 0 to 0.0500%;
As: 0 to 0.0500%;
Mg: 0 to 0.0500%;
Ca: 0 to 0.0500%;
Y: 0 to 0.0500%;
30 Zr: 0 to 0.0500%;
La: 0 to 0.0500%;
Ce: 0 to 0.0500%; and
a balance of Fe and impurities,
wherein
35 the steel sheet comprises, by area ratio:
a total of ferrite and bainite: 10.0 to 90.0%;
48
a total of martensite and tempered martensite: 5.0 to 80.0%; and
a total of pearlite and retained austenite: 0 to 15.0%,
integration degree of (111)<112> orientation of ferrite is 3.0 or higher, and
integration degree of (252)<2-11> orientation of martensite and tempered martensite is 5.0
5 or lower.
[Claim 2]
The steel sheet according to claim 1, comprising one or more of:
Co: 0.0001 to 0.5000%;
10 Ni: 0.0001 to 0.5000%;
Mo: 0.0001 to 0.5000%;
Cr: 0.0001 to 1.0000%;
O: 0.0001 to 0.0200%;
Ti: 0.0001 to 0.5000%;
15 B: 0.0001 to 0.0100%;
Nb: 0.0001 to 0.5000%;
V: 0.0001 to 0.5000%;
Cu: 0.0001 to 0.5000%;
W: 0.0001 to 0.1000%;
20 Ta: 0.0001 to 0.1000%;
Sn: 0.0001 to 0.0500%;
Sb: 0.0001 to 0.0500%;
As: 0.0001 to 0.0500%;
Mg: 0.0001 to 0.0500%;
25 Ca: 0.0001 to 0.0500%;
Y: 0.0001 to 0.0500%;
Zr: 0.0001 to 0.0500%;
La: 0.0001 to 0.0500%; and
Ce: 0.0001 to 0.0500%.
30
[Claim 3]
A method of producing a steel sheet, the method comprising:
a casting step of continuously casting a molten steel having the chemical composition
according to claim 1 or 2 to form a slab, wherein 5 to 40% rolling reduction is performed at a
35 temperature of 800C to lower than 1,200C in a period after the continuous casting and before
cooling to room temperature;
49
a hot rolling step which includes hot rolling the slab and in which a finishing temperature of
the hot rolling is 650 to 950C;
a step of coiling the thus obtained hot-rolled steel sheet at a coiling temperature of 400 to
700C;
a step of retaining the thus coiled hot 5 -rolled steel sheet as is wherein the thus coiled
hot-rolled steel sheet is not cooled to room temperature before retaining and is retained in a
temperature range of (coiling start temperature + 20C to 100C) for 5 to 300 minutes;
a cold rolling step of cold rolling the hot-rolled steel sheet at a rolling reduction ratio of
10.0 to 90.0%; and
10 an annealing step of annealing the thus obtained cold-rolled steel sheet in a temperature
range of 700 to 900C.