[Title of the Invention] MIDDLEIHIGH CARBON STEEL SHEET AND
METHOD FOR MANUFACTURING SAME
[Technical Field of the Invention]
[OOOI]
The present invention relates to a middlelhigh carbon steel sheet exhibiting
excellent reduction in area during shaping at a high strain rate and a method for
manufacturing the same.
The present application claims priority on the basis of Japanese Patent Application No.
2014-045689, filed on March 7,2014, the content of which is incorporated herein.
[Related Art]
[0002]
Middlehigh carbon steel sheets are used as materials for drive system
components such as chains, gears, and clutches in vehicles, saws, blades, and the like.
Materials obtained by shaping steel strips of middlelhigh carbon steel or steel sheets
cut out from the steel strips into predetermined shapes are shaped into component
shapes by deformation processing such as deep drawing, hole expanding, thickening,
or thinning. In cold forging in which each of processes is individually carried out or
processes of multiple kinds are carried out at the same time, materials are partially
shaped at a high strain rate of approximately 10 Isec, and, for steel sheets that are used
as materials, there is a demand for excellent formability, that is, excellent reduction in
area even during distortion at a high strain rate.
[0003]
Thus far, there have been a number of proposals regarding techniques that
improve the reduction in area of middlehigh carbon steel sheets (for example, refer to
Patent Documents 1 to 6).
[00041
For example, Patent Document 1 discloses an invention of a method for
manufacturing a middlelhigh carbon steel sheet having excellent deep drawability, in
which finishing rolling is carried out on a hot-rolled steel sheet or an annealed steel
sheet containing C: 0.20% by mass to 0.90% by mass using a work roller having a
surface roughness Ra in a range of 0.20 pm to 1.50 pm in at least a final rolling path
under conditions ofthe total rolling reduction being set in a range of 20% to 70%, and
then finishing annealing is carried out. However, the technique disclosed by Patent
Document 1 is a technique in which reduction in area is increased by improving the
roughness of the steel sheet surface, but is not a technique in which reduction in area is
increased by improving material quality by the control of the structure forms of steel
products and thus does not always provide the desired effects of the invention.
Furthermore, I'atent Document 2 discloses an invention of a high-roughness
high carbon steel sheet having excellent workability, including C:0.6% by mass to
1.3% by mass, Si: 0.5% by mass or less, Mn: 0.2% by mass to 1.0% by mass, P: 0.02%
by mass or less, and S: 0.01% by mass or less with a remainder substantially having a
composition of Fe, in which, by adjustment of hot-rolling conditions, cold-rolling
conditions, and annealing conditions, the maximum length of carbides is set to be
equal to or shorter than 5.0 pm, the carbide spheroidizing ratio is set to be equal to or
higher than 90%, the volume of spherical carbides having a grain size of equal to or
larger than 1.0 pm is set to be equal to or higher than 20% of the total spherical carbide
volume, and the high carbon steel sheet is made up of carbides and equiaxial ferrite.
[0006]
Patent Document 3 discloses an invention of middlethigh carbon steel
exhibiting excellent reduction in size, in which the C content is in a range of 0.10% by
mass to 0.90% by mass, and a structure in which carbides are dispersed in ferrite so
that a ferrite intergranular abundance (F value) of the carbides reaches equal to or
higher than 30% is formed.
[0007]
Patent Document 4 discloses an invention of a high carbon cold-rolled steel
strip which is slightly anisotropic in a deep drawn surface, having a steel composition
of C: 0.25% to 0.75%, in which the average grain size of carbides in steel is equal to or
larger than 0.5 puthen sp,h eroidizing ratio is equal to or higher than 90%, and a texture
satisfies an expression "(222)/(200)~6-8.0xC (%)".
[OOOS]
Patent Document 5 discloses an invention of a high carbon steel strip which
has favorable deep drawability and, furthermore, is capable of imparting high strength
or excellent wear resistance, in which the C content is in a range of 0.20% by mass to
0.70% by mass, and equal to or higher than 50% by area of cementite in the steel is
graphitized.
[0009]
Patent Document 6 discloses a technique of a method for manufacturing a
high carbon cold-rolled steel sheet having excellent formability, in which high carbon
steel containing C: 0.1% to 0.65%, Si: 0.01% to 0.3%, Mn: 0.4% to 2%, sol. Al: 0.01%
to 0.1%, N: 0.002% to 0.008%, R: 0.0005% to 0.005%, Cr: 0 to 0.5, and Mo: 0 to 0.1
is hot-rolled, is coiled at 300°C to 520°C, is box-annealed at 650°C to (Acl-lO)'C, is
cold-rolled at a rolling reduction in a range of 40% to 80%, and is box-annealed at
650°C to (Acl-1O)OC.
[OO 101
However, none of the above-described patent documents discloses anything
about knowledge and techniques that suppress the cracking of cementite in steel
products, which occurs during shaping at a high strain rate, and a decrease in reduction
in area caused by the growth and joining of voids initiated due to the initiation of
cracks.
[Prior Art Document]
[Patent Document]
[OOll]
[Patent Document 11 Japanese Unexamined Patent Application, First Publication No.
2003-293042
[Patent Document 21 Japanese Unexamined Patent Application, First Publication No.
2003-147485
[Patent Document 31 Japanese Unexamined Patent Application, First Publication No.
2002-155339
[Patcnt Document 41 Japanese Unexamined Patent Application, First Publication No.
2000-328 172
[Patent Document 51 Japanese Unexamined Patent Application, First Publication No.
H06-108158
[Patent Document 61 Japanese Unexamined Patent Application, First Publication No.
H11-61272
[Disclosure of the Invention]
[Problems to be Solved by the Invention]
[OO 121
The present invention has been made in consideration of the above-described
circumstances, and an object of the present invention is to provide a middlehigh
carbon steel sheet exhibiting excellent reduction in area during shaping at a high strain
rate and a method for manufacturing the same.
[Means for Solving the Problem]
[0013]
The present inventors carried out intensive studies regarding methods for
achieving the above-described object. As a result, the present inventors found that
cracks (voids) forming at carbides during distortion propagate and join together, and
thus reduction in area is decreased during distortion at a high strain rate. Furthermore,
the present inventors found that cracks forming at carbides are initiated from crystal
interfaces present in carbide particles which have been considered as a single particle
in the related art. The present inventors found that, when the amount of crystal
interfaces in carbide particles is decreased, it is possible to obtain a middlehigh carbon
steel sheet which exhibits excellent reduction in area even during distortion at a high
strain rate and, furthermore, exhibits excellent formability in cold forging in which
deformation processing such as deep drawing, hole expanding, thickening, or thinning
is carried out or multiple kinds thereof are carried out at the same time.
[0014]
In addition, the present inventors repeated a variety of studies and thus found
that it is~dificultto manufacture steel sheets having the above-described characteristics
in a case in which efforts are made to separately find appropriate hot-rolling conditions,
annealing conditions, and the likei and the steel sheets can only be manufactured by
achicving optimization by so-called collective processing such as hot-rolling and
annealing processing, and completed the present invention.
[0015]
The outline of the present invention is as described below.
[OO 161
(1) An aspect of the present invention provides a middlehigh carbon steel
sheet, in which composition thereof contains, by mass%, C: 0.10% to 1.50%, Si:
0.01% to 1.00%, Mn: 0.01% to 3.00%, P: 0.0001% to 0.1000%, and S: 0.0001% to
0.1000%, and a remainder consisting of Fe and impurities, in which the steel sheet has
a structure in which a total volume percentage of a martensite, a bainite, a pearlite, and
a residual austenite is equal to or lower than 5.0%, and a remainder thereof is a ferrite
and carbides, in which the spheroidizing ratio of carbide particles is 70% to 99%, and
in which a proportion of a number of the carbide particles including a crystal interface
at which an orientation difference is equal to or greater than 5" in the carbide particles
is equal to or lower than 20% of a total number of the carbide particles.
[0017]
(2) The middleihigh carbon steel sheet according to (I), in which thc
composition of the steel sheet may further include, by mass%, one or more selected
from the group consisting of Al: 0.001% to 0.500%, N: 0.0001% to 0.0500%, 0:
0.0001% to 0.0500%, Cr: 0.001% to 2.00%, Mo: 0.001% to 2.000%, Ni: 0.001% to
2.00%, Cu: 0.001% to 1.000%, Nb: 0.001% to 1.000%, V: 0.001% to 1.000%, Ti:
0.001% to 1.000%, B: 0.0001% to 0.0500%, W: 0.001% to 1.000%, Ta: 0.001% to
1.000%, Sn: 0.001% to 0. 020%, Sb: 0.001% to 0.020%, As: 0.001% to 0.020%, Mg:
0.0001% to 0.0200%, Ca: 0.001% to 0.020%, Y: 0 .001% to 0.020%, Zr: 0.001% to
0.020%, La: 0.001% to 0.020%, and Ce: 0.001% to 0.020%.
[OOlS]
(3) Another aspect of the prescnt invention provides a method for
manufacturing a middlehigh carbon steel sheet, in which, when a billet having the
composition according to (1) or (2) is directly hot-rolled or temporary cooled, heated,
and hot-rolled, finish hot-rolling is completed in a temperature region of 600°C to
1,00O0C, the hot-rolled steel sheet coiled at 350°C to 700' is box-annealed, coldrolling
of 10% to 80% is carried out, and then cold-rolled-sheet-annealing is carried
out at an annealing temperature of 650°C to 780°C for a retention time of 30 to 1,800
seconds in a continuous annealing line.
[Effects of the Invention]
[0019]
According to the present invention, it is possible to provide a middlehigh
carbon steel sheet exhibiting excellent reduction in area during shaping at a high strain
rate and a method for manufacturing the same.
[Brief Description of the Drawing(s)]
[0020]
FIG. 1 is a view showing a shape of a test specimen used for measuring
reduction in area at a high strain rate.
FIG. 2 is a view showing an appearance in which cracks are initiated from
crystal interfaces present in carbide particles during distortion
FIG. 3 is a view showing a relationship between a proportion of the number of
carbide particles including a crystal interface and reduction in area during a tensile test
at a high strain rate.
[Embodiment(s) of the Invention]
[0021]
Hereinafter, the present embodiment will be described in detail.
[0022]
First, the reasons for limiting the chemical composition of a stecl shect
according to the present cmbodirnent will be described. Here, regarding the
composition, "%" represents "mass%.
[0023]
(C: 0.10% to 1.50%)
C is an element that increases the strength of steel by a heat treatment of
quenching. Middlelhigh carbon steel sheets ensure strength or toughness necessary
for components by heat treatments such as quenching and quenching and tempering
which are carried out after shaping and before the use of the steel sheets as materials
for drive system components such as chains, gears, and clutches in vehicles, saws,
blades, and the like. When the C content is lower than 0.10%, the strength cannot be
increased by quenching, and thus the lower S i t of the C content is set to 0.10%. On
the other hand, when the C content exceeds 1.50%, after cold-rolling and annealing,
the proportion of the number of carbides including a crystal interface in the particle
increases, and reduction in area is decreased at a high strain rate, and thus the upper
limit of the C content is set to 1.50%. More preferably, the C content is in a range of
0.15% to 1.30%.
[0024]
(Si: 0.01% to 1.00%)
Si is an element that acts as a deoxidizing agent and suppresses the coarsening
and joining of carbide particles during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing. In a process of the Ostwald growth of carbide particles during coldrolled-
sheet-annealing, when two or more particles that are adjacent to each other
come into contact with each other, crystal interfaces are introduced into thc carbide
particles. During the distortion ofthe steel sheet, the crystal interfaces in the carbide
particles serve as the starting points of cracks. In order to suppress the abovedescribed
phenomenon, it is necessary to decrease the growth rate ol' carbides during
hot-rolled-sheet-annealing and cold-rolled-sheet-annealing. One of the typical
elements that decrease the growth rate of carbides during hot-rolled-sheet-annealing
and cold-rolled-sheet-annealing is Si. When the Si content is lower than 0.01%, the
above-described effects camlot be obtained, and thus the lower limit of the Si content
is set to 0.01%. On the other hand, when the Si content exceeds 1.00%, ferrite
becomes prone to cleavage fracture, and reduction in area is decreased at a high strain
rate, and thus the upper limit of the Si content is set to 1.00%. The Si content is more
preferably 0.05% to 0.80% and still more preferably 0.08% to 0.50%.
[0025]
(Mn: 0.01% to 3.00%)
Mn is, similar to Si, an element that suppresses the coarsening and joining of
carbide particles during hot-rolled-sheet-annealing and cold-rolled-sheet-annealing.
When the Mn content is lower than 0.01%, the above-described effects cannot be
obtained, and thus the lower limit of the Mn content is set to 0.01%. On the other
hand, when the Mn content exceeds 3.00%, it becomes difficult for carbides to
spheroidize during hot-rolled-sheet-annealing and cold-rolled-sheet-annealing, cracks
are initiated from needle-like carbides as starting points during distortion at a high
strain rate, and reduction in area decreases. Therefore, the npper limit of the Mn
content is set to 3.00%. The Mn content is more preferably 0.30% to 2.50% and still
more preferably 0.50% to 1.50%.
[0026]
(P: 0.0001% to 0.1000%)
P is an impurity element that embrittles grain boundaries of ferrite. The P
content is preferably lower; however, in a case in which steel is highly purified by
setting the P content to be lower than 0.0001% in refining, a time necessary for
refining becomes long, and manufacturing costs are significantly increased, and thus
the lower limit of the P content is set to 0.0001%. On the other hand, when the P
content exceeds 0.1000%, cracks are significantly initiated from grain boundaries of
ferrite during distortion at a high strain rate, and reduction in area is significantly
decreased, and thus the upper limit of the P content is set to 0.1000%. The P content
is more preferably 0.0010% to 0.0500% and still more preferably 0.0020% to 0.0300%.
[0027]
(S: 0.0001% to 0.1000%)
S is an impurity element that forms non-metallic inclusions such as MnS, and
non-metallic inclusions act as starting points for the initiation of cracks during
distortion at a high strain rate, and thus the S content is preferably lower. However, a
decrease in the S content to lower than 0.0001% leads to a significant increase in
refining costs, and thus the lower limit of the S content is set to 0.0001%. On the
other hand, when higher than 0.1000% of S is included, reduction in area is
significantly decreased, and thus the upper limit of the S content is set to equal to or
lower than 0.1000%. The S content is more preferably 0.0003% to 0.0300%.
[0028]
In the present embodiment, the above-described composition is the base
elements of the steel sheet; however, it is also possible to further, optionally, add one
or two or more selected from the elements described below in order to improve the
mechanical characteristics of the steel sheet. Here, the elements described below do
not need to he essentially included, and thus the lower limit values of the elements
described below are 0%.
[0029]
(Al: Preferably 0.001% to 0.500%)
A1 is an element that serves as a deoxidizing agent of steel. When the A1
content is lower than 0.001%, effects ofthe inclusion of A1 cannot be sufficiently
obtained, and thus the lower limit of the A1 content may be set to 0.001%. On the
other hand, when the A1 content exceeds 0.500%, grain boundaries of ferrite are
embrittled, and reduction in area during distortion at a high strain rate is decreased.
Therefore, the upper limit of the A1 content may be set to 0.500%. The A1 content is
more preferably 0.005% to 0.300% and still more preferably 0.010% to 0.100%.
[0030]
(N: Preferably 0.0001% to 0.0500%)
N is an element that accelerates the bainite transformation of steel. In
addition, N causes ferrite to embrittle when a large amount of N is included. The N
content is preferably lower, but a decrease in the N content to lower than 0.0001%
leads to an increase in refining costs, and thus the lower limit of the N content may be
set to 0.0001%. On the other hand, when the N content exceeds 0.0500%, the
cracking of ferrite is caused during distortion at a high strain rate, and thus the upper
limit of the N content may be set to 0.0500%. The N content is more preferably
0.0010% to 0.0250% and still more preferably n 0.0020% to 0.0100%.
[003 11
(0: Preferably 0.0001% to 0.0500%)
0 is an element that accelerates the formation of coarse oxides in steel when a
large amount of 0 is included, and thus the 0 content is preferably lower. However,
a decrease in thc 0 content to lower than 0.0001 % leads to an increase in refining costs.
and thus the lower limit of the 0 content may be set to 0.0001%. On the other hand,
when the 0 content exceeds 0.0500%, coarse oxides are formed in steel, and cracks are
initiated from the coarse oxides as starting points during distortion at a high strain rate,
and thus the upper limit of the 0 content may be set to 0.0500%. The 0 content is
more preferably 0.0005% to 0.0250% and still more preferably 0.0010% to 0.0100%.
[0032]
(Cr: Preferably 0.001% to 2.000%)
Cr is an element that, similar to Si and Mn, suppresses the coarsening and
joining of carbide particles during hot-rolled-sheet-annealing and cold-rolled-sheetannealing.
However, when the Cr content is lower than 0.001%, the above-described
effect cannot be obtained, and thus the lower limit of the Cr content may he set to
0.001%. On the other hand, when the Cr content exceeds 2.000%, it becomes
difficult for carbides to spheroidize during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing, cracks are initiated from needle-like carbides as starting points during
distortion at a high strain rate, and reduction in area is decreased, and thus the upper
limit of the Cr content may he set to 2.000%. The Cr content is more preferably
0.005% to 1.500% and still more preferably 0.010% to 1.300%.
[0033]
(Mo: Preferably 0.001% to 2.000%)
Mo is an clement that, similar to Si, Mn, and Cr, suppresses the coarsening
and joining of carbide particles during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing. When the Mo content is lower than 0.001%, the above-described
effect cannot be obtained, and thus the lower limit ofthe Mo content may be set to
0.001%. On the other hand, when the Mo content exceeds 2.000%, it becomes
difficult for carbides to spheroidize during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing, cracks are initiated from needle-like carbides as starting points during
distortion at a high strain rate, and reduction in area is decreased, and thus the upper
limit of the Mo content may be set to 2.000%. The Mo content is more preferably
0.005% to 1.900% and still more preferably 0.008% to 0.800%.
[0034]
mi: Preferably 0.001% to 2.000%)
Ni is an elemcnt effective for improving the toughness of components and
improving hardenability. In order to effectively exhibit the above-described effect,
equal to or higher than 0.001% of Ni is preferably included. On the other hand, when
the Ni content exceeds 2.000%, it becomes difficult for carbides to spheroidize during
hot-rolled-sheet-annealing and cold-rolled-sheet-annealing, cracks are initiated from
needle-like carbides as starting points during distortion at a high strain rate, and
reduction in area is decreased, and thus the upper limit of the Ni content may he set to
2.000%. The Ni content is more preferably 0.005% to 1.500% and still more
preferably 0.005% to 0.700%.
[0035]
(Cu: Preferably 0.001% to 1.000%)
Cu is an element that increases the strengths of steel products by forming fine
precipitates. In order to effectively exhibit the effect of an increase in strength, equal
to or higher than 0.001% of Cu is preferably included. On the other hand, when the
Cu content exceeds 1.00%, it becomes difficult for carbides to spheroidize during hotrolled-
sheet-annealing and cold-rolled-sheet-annealing, cracks are initiated from
needle-like carbides as starting points during distortion at a high strain rate, and
reduction in area is decreased, and tbt~sth e upper limit of the Cu content may be set to
1.00%. The Cu content is more preferably 0.003% to 0.500% and still more
preferably 0.005% to 0.200%.
[0036]
(Nb: Preferably 0.001% to 1.000%)
Nb is an element that forms carbonitrides and suppresses the coarsening and
joining of carbide particles during hot-rolled-sheet-annealing and cold-rolled-sheetannealing.
However, when the Nb content is lower than 0.001%, the above-described
effect cannot be obtained, and thus the lower limit of the Nb content may be set to
0.001%. On the other hand, when the Nb content exceeds 1.000%, it becomes
difficult for carbides to spheroidize during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing, cracks are initiated from needle-like carbides as starting points during
distortion at a high strain rate, and reduction in area is decreased, and thus the upper
limit of the Nb content may be set to 1.000%. The Nb content is more preferably
0.005% to 0.600% and still more preferably 0.008% to 0.200%.
[0037]
(V: Preferably 0.001% to 1.000%)
V is also an clement that, similar to Nb, forms carbonitrides and suppresses
the coarsening and joining of carbide particles during hot-rolled-sheet-annealing and
cold-rolled-sheet-annealing. When the V content is lower than 0.001%, the abovedescribed
effect cannot be obtained, and thus the lower limit of the V content may be
set to 0.001%. On the other hand, when the V content exceeds 1.000%, it becomes
difficult for carbides to spheroidize during hot-rolled-sheet-annealing and cold-rolledsheet-
annealing, cracks are initiated from needle-like carbides as starting points during
distortion at a high strain rate, and reduction in area is decreased, and thus the upper
limit of the V contcnt may be set to 1.000%. The V content is more preferably
0.001% 0.750% and still more preferably 0.001% to 0.250%.
[003 81
(Ti: Preferably 0.001% to 1.000%)
Ti is also an element that, similar to Nb and V, forms carbonitrides and
suppresses the coarsening and joining of carbide particles during hot-rolled-sheetannealing
and cold-rolled-sheet-annealing. When the Ti content is lower than
0.001%, the above-described effect cannot be obtained, and thus the lower limit of the
Ti content may be set to 0.001%. On the other hand, when the Ti content exceeds
1.000%, it becomes difficult for carbides to spheroidize during hot-rolled-sheetannealing
and cold-rolled-sheet-annealing, cracks are initiated from needle-like
carbides as starting points during distortion at a high strain rate, and reduction in area
is decreased, and thus the upper limit of the Ti content may be set to 1.000%. The Ti
content is more preferably 0.001% to 0.500% and still more preferably 0.003% to
0.150%.
[0039]
(B: Preferably 0.0001% to 0.0500%)
B is an element that improves hardenability during a heat treatment of
components. When the B content is lower than 0.0001%, the above-described effect
cannot be obtained, and thus the lower limit of the B content may be set to 0.0001%.
When the B content exceeds 0.0500%, coarse Fe-B-C compo~mdsa re generated and
serve as starting points during distortion at a high strain rate, and reduction in area is
decreased, and thus the upper limit of the B content may be set to 0.0500%. The B
content is more preferably 0.0005% to 0.0300% and still more preferably 0.0010% to
0.0100%.
[0040]
(W: Preferably 0.001% to 1.000%)
W is also an element that, similar to Nb, V, and Ti, forms carbonitrides and
suppresses the coarsening and joining of carbide particles during hot-rolled-sheetannealing
and cold-rolled-sheet-annealing. When the W content is lower than
0.001%, the above-described effect cannot be obtained, and thus the lower limit of the
W content may be set to 0.001%. On the other hand, when the W content exceeds
1.000%, it becomes difficult for carbides to spheroidize during hot-rolled-sheetannealing
and cold-rolled-sheet-annealing, cracks are initiated from needle-like
carbides as starting points of cracks during distortion at a high strain rate, and
reduction in area is decreased, and thus the upper limit of the W content may be set to
1.000%. The W content is more preferably 0.001% to 0.450% and still more
preferably 0.001% to 0.160%.
[0041]
(Ta: Preferably 0.001% to 1.000%)
Ta is also an element that, similar to Nb, V, Ti, and W, forms carbonitrides
and suppresses the coarsening and joining of carbide particles during hot-rollcd-sheetannealing
and cold-rolled-sheet-annealing. When the Ta content is lower than
0.001%, the above-described effect cannot be obtained, and thus the lower limit of the
Ta content may be set to 0.001%. On the other hand, when the Ta content exceeds
1.000%, it becomes difficult for carbides to spheroidize during hot-rolled-sheetannealing
and cold-rolled-sheet-annealing, cracks are initiated from needle-like
carbides as starting points during distortion at a high strain rate, and reduction in area
is decreased, and thus the upper limit of the Ta content may be set to 1.000%. The Ta
content is more preferably 0.001% to 0.750% and still more preferably 0.001% to
0.150%.
100421
(Sn: Preferably 0.001% to 0.020%)
Sn is an element included in steel in a case in which scraps are used as a steel
raw material, and the Sn content is preferably lower. In a case in which the Sn
content is decreased to lower than 0.001%, refining costs are increased, and thus the
lower limit of the Sn content may be set to 0.001%. In addition, in a case in which
the Sn content exceeds 0.020%, ferrite embrittles, and reduction in area is decreased
during distortion at a high strain rate, and thus the upper limit of the Sn content may be
set to 0.020%. The Sn content is more preferably 0.001% to 0.015% and still more
preferably 0.001% to 0.0 10%.
[0043]
(Sb: Preferably 0.001% to 0.020%)
Sb is, similar to Sb, an element included in steel in a case in which scraps are
used as a steel raw material, and the Sb content is preferably lower. In a case in
which the Sb content is decreased to lower than 0.001%, refining costs are increased,
and thus the lower limit of the Sb content may be set to 0.001%. In addition, in a case
in which the Sb content exceeds 0.020%, ferrite embrittles, and reduction in area is
decreased during distortion at a high strain rate, and thus the upper limit of the Sb
content may be set to 0.020%. The Sb content is more preferably 0.001% to 0.015%
and still more preferably 0.001% to 0.01 1%.
[0044]
(As: Preferably 0.001% to 0.020%)
As is, similar to Sn and Sb, an element included in steel in a case in which
scraps are used as a steel raw material, and the As content is preferably lower. In a
case in which the As content is decreased to lower than 0.001%, refining costs are
increased, and thus the lower limit of the As content may be set to 0.001%. In
addition, in a case in which the As content exceeds 0.020%, ferrite embrittles, and
reduction in area is decreased during distortion at a high strain rate, and thus the upper
limit of the As content may be set to 0.020%. The As content is more preferably
0.001% to 0.015% and still more preferably 0.001% to 0.007%.
[0045]
(Mg: Preferably 0.0001% to 0.0200%)
Mg is an element capable of controlling the form of sulfides even when the
content thereof is low and can be included as necessary. When the Mg content is
lower than 0.0001%, the above-described effect cannot be obtained, and thus the lower
limit of the Mg content may be set to 0.0001%. On the other hand, in a case in which
Mg is excessively included, grain boundaries of ferrite are embrittled, and reduction in
area during distortion at a high strain rate is decreased, and thus the upper limit of the
Mg content may be set to 0.0200%. The Mg content is more preferably 0.0001% to
0.0150% and still more preferably 0.0001% to 0.0075%.
200461
(Ca: Preferably 0.001% to 0.020%)
Ca is, similar to Mg, an element capable of controlling the form of sulfides
even when the content thereof is low and can be included as necessary. When the Ca
content is lower than 0.001%, the above-described effect cannot be obtained, and thus
the lower limit of the Ca content may be set to 0.001%. On the other hand, in a case
in which Ca is excessively included, grain boundaries of ferrite are embrittled, and
reduction in area during distortion at a high strain rate is decreased, and thus the upper
limit of the Ca content may be set to 0.020%. The Ca content is more preferably
0.001% to 0.015% and still more preferably 0.001% to 0.010%.
[0047]
(Y: Preferably 0.001% to 0.020%)
Y is, similar to Mg and Ca, an element capable of controlling the form of
sulfides even when the content thereof is low and can be included as necessary.
When the Y content is lower than 0.001%, the above-described effect cannot be
obtained, and thus the lower limit of the Y content may be set to 0.001%. On the
other hand, in a case in which Y is excessively included, grain boundaries of ferrite are
embrittled, and reduction in area during distortion at a high strain rate is decreased, and
thus the upper limit of the Y content may be set to 0.020%. The Y content is more
preferably 0.001% to 0.015% and still more preferably 0.001% to 0.009%.
[0048]
(Zr: Preferably 0.001% to 0.020%)
Zr is, similar to Mg, Ca, and Y, an element capable of controlling the form of
sulfides even when the content thereof is low and can be included as necessaly.
When the Zr content is lower than 0.001%, the above-described effect cannot be
obtained, and thus the lower limit of the Zr content may be set to 0.001%. On the
other hand, in a case in which Zr is excessively included, grain boundaries of ferrite
are embrittled, and reduction in area during distortion at a high strain rate is decreased,
and thus the upper limit of the Zr content may be set to 0.020%. The Zr content is
more preferably equal to or lower than 0.01 5% and still more preferably equal to or
lower than 0.010%.
[0049]
(La: Preferably 0.001% to 0.020%)
La is, similar to Mg, Ca, Y, and Zr, an element capable of controlling the form
of sulfides even when the content thereof is low and can be included as necessary.
When the La content is lower than 0.001%, the above-described effect cannot be
obtained, and thus the lower limit of the La content may be set to 0.001%. On the
other hand, in a case in which La is excessively included, grain boundaries of ferrite
are embrittled, and reduction in area during distortion at a high strain rate is decreased,
and thus the upper limit of the La content may be set to 0.020%. The La content is
more preferably 0.001% to 0.015% and still more preferably 0.001% to 0.010%.
[0050]
(Ce: Preferably 0.001% to 0.020%)
Ce is, similar to Mg, Ca, Y, Zr, and La, an element capable of controlling the
form of sulfides even when the content thereof is low and can be included as necessary.
When the Ce content is lower than 0.001%, the above-described effect cannot be
obtained, and thus the lower limit of the Ce content may be set to 0.001%. On the
other hand, in a case in which Ce is excessively included, grain boundaries of ferrite
are embrittled, and reduclion in area during distortion at a high strain rate is decreased,
and thus the upper limit of the Ce content may be set to 0.020%. The Ce content is
more preferably 0.001% to 0.015% and still more preferably 0.001% to 0.010%.
[OOSl]
Meanwhile, in the steel sheet according to the present embodiment, the
remainder of the composition described above is Fe and impurities.
lo0521
The steel sheet according to the present embodiment does not only have the
above-described composition but is also subjected to optimal hot-rolling and annealing
and thus the steel sheet has a structure in which ferrite and carbides are main bodies,
the total volume percentage of martensite, bainite, pearlite, and residual austenite is
equal to or lower than 5.0%, the spheroidizing ratio of carbide particles is 70% to 99%,
and the proportion ofthe number of the carbide particles including a crystal interface at
which an orientation difference is equal to or greater than 5" in the carbide particles is
equal to or lower than 20% of the total number of the carbide particles. Due to these
characteristics, it is possible to obtain steel sheets having excellent formability when
deformation processing such as reduction in area, hole expanding, thickening, or
thinning or cold forging in which the above-described processes are combined together
is carried out at a high strain rate. This is new knowledge that the present inventors
found.
[0053]
Steel according to the present embodiment has a structure of substantially
ferrite and carbides. Meanwhile, carbides refer to not only cementite (Fe3C) which is
a compound of iron and carbon but also a compound in which Fe atoms in cementite
are substituted with alloy elements such as Mn and Cr and alloy carbides (M23C6, MsC,
MC; here, M represents Fe and other alloy elements). Martensite, bainite, pcarlite,
and residual austenite are preferably not included in the structure, and, in a case in
which they are included, the total volume percentage is set to equal to or lower than
5.0%. The lower limit of the total amount of martensite, bainite, pearlite, and residual
austenite is not regulated. In a case in which no structures thereof are detected in a
structure observation at a magnification of 3,000 times using a scanning electron
microscope, which will be described below, the total amount of martensite, bainite,
pearlite, and residual austenite is considered as 0.0% by volume, and thus the lower
limit of the total amount of martcnsite, bainite, pearlite, and residual austenite may be
set to 0.0%.
[0054]
The reasons for limiting the total amount of martensite, bainite, pearlite, and
residual austenite will be described. Martensite, bainite, pearlite, and residual
austenite which are the regulation subjects in the present embodiment are structures
generated from austenite in a process in which the steel sheet is heated to a two-phase
region of ferrite and austenite during cold-rolled-sheet-annealing and then is cooled to
room temperature. Therefore, martensite, bainite, and pearlite are located in grain
boundaries of ferrite, and residual austenite is present in lath interfaces or block
boundaries between martensite and bainite. First, when austenite transforms to
martensite, bainite, or pearlite, the volume expands, and thus stress remains in grain
boundaries of ferrite. Stress locally remaining in the grain boundaries of ferrite
accelerate the initiation of voids in the vicinities of the grain boundaries during
distortion of the steel sheet due to stress loading, and thus stress remaining the grain
boundaries of ferrite leads to a decrease in reduction in area during distortion at a high
strain rate. In addition, residual austenite turns into martensite during the distortion
of the steel sheet by processing-induced transformation caused therein, and thus an
increase in stress in the ferrite grain boundaries is further increased, and a decrease in
reduction in area is promoted. For the above-described reasons, in order to improve
reduction in area during distortion at a high strain rate, it is preferable to set the
structure of the steel sheet to a structure of substantially ferrite and carbides and to
include no martensite, bainite, pearlite, and residual austenite in the structure, and, in a
case in which they are included, it becomes essential to set the total volume percentage
of martensite, bainite, pearlite, and residual austenite to equal to or lower than 5.0%.
Furthermore, in a case in which pearlitic transformation is caused, the proportion of
needle-like carbides also increases. The influences of needle-like carbides will be
described below. Meanwhile, in carbides, phase transformation does not occur, and
stress does not accumulate between carbides and base metal, and thus it is possible to
limit a decrease in reduction in area.
LO0551
Next, the reasons for setting the spheroidizing ratio of carbides to 70% to 99%
will be described. When the spheroidizing ratio of carbides is lower than 70%, stress
accumulates at needle-like carbides, carbides crack, thus, voids are initiated, and voids
joined together form a broken surface, and thus reduction in area during distortion at a
high strain rate is decreased. Therefore, the lower limit of the spheroidizing ratio of
carbides is set to 70%. Meanwhile, the spheroidizing ratio is desirably higher;
however, in order to control the spheroidizing ratio to be 100%, it is necessary to carry
out annealing for an extremely long period of time, which leads to an increase in
manufacturing costs, and thus the upper limit of the spheroidizing ratio is desirably
lower than 100% and is set to equal to or lower than 99%.
[0056]
Furthermore, the reasons for setting the proportion of the number of carbide
particles including a crystal interface at which a crystal orientation difference is equal
to or greater than 5" in carbide particles to be equal to or lower than 20% of the total
:, number of carbide particles will be described. Cracking of carbides during distortion
I!
1, mainly initiates from crystal interfaces at which a crystal orientation difference is equal
1,
:I to or greater than 5" which are present in the carbides which have been considered as a
I
single particle in the related art. During distortion at a high strain rate, voids are
initiated due to the cracking of carbides at crystal interfaces, the voids join together,
- 23 -
and a broken surface is formed, whereby reduction in area is decreased. The
proportion of carbides including a crystal interface at which a crystal orientation
difference is equal to or greater than 5" is preferably lower; however, in order to
control the proportion of the number of carbides including a crystal interface at which
a crystal orientation difference is equal to or greater than 5" to he lower than 0.1% of
the total number of carbide particles, collective quality design management becomes
essential in continuous forging, hot-rolling, hot-rolled-sheet-annealing, cold-rolling,
and cold-rolled-sheet-annealing, and the yield is decreased, and thus the lower limit of
the proportion of the number of carbides including a crystal interface at which a crystal
orientation difference is equal to or greater than 5' of the total number of casbide
particles is preferably set to 0.1% and more preferably set to 0.2%. In addition, in a
case in which the proportion of the number of carbides including a crystal interface at
which an orientation difference is equal to or greater than 5" in the total number of
carbide particles exceeds 20%, reduction in area is significantly decreased during
distortion at a high strain rate, and thus the upper limit of the proportion of the number
thereof is set to 20% and is more preferably 15% and still more preferably 10%.
[0057]
Subsequently, a method for observing and measuring the stmcture regulated
above will be described.
[OOS 81
Ferrite, carbides, martcnsite, bainite, and pearlite are observed using a
scanning electron microscope. Before observation, samples for structural observation
are wet-polished using Emery paper and are polished using diamond abrasive grains
having an average particle size of 1 bm, thereby finishing the observed sections to be
mirror-like surfaces. Next, the observed ~ectionsa re etched using a 3% nitric acidalcohol
solution. Regarding the observation magnification, a magnification at which
determination of the respective structures of ferrite, carbides, martensite, bainite, and
pearlite becomes possible is selected in a range of 1,000 times to 10,000 times. In the
present embodiment, a magnification of 3,000 times was selected. At the selected
magnification, 30 ymx40 ym visual fields randomly taken from a quarter thickness
layer are captured 16 times. The volume percentages of the respective structures are
obtained using a point count method. On captured structural photographs, grid lines
are drawn vertically and horizontally at intervals of 2 ym, the numbers of the structures
at the intersections of the grid lines are respectively counted, and the proportions of the
respective structures per the captured photograph are measured from the proportions of
the numbers of the respective structures. After that, the average values of the
measurement results of the proportions of the respective structures according to all of
the 16 structural photographs are obtained as the volume percentages of the structures
in the respective samples.
[0059]
Meanwhile, martensite and bainite are differentiated on the basis of the
presence or absence of fine carbides in the structure. A structure which is mainly
located on a grain boundary of ferrite and does not include carbides is martensite, and a
structure including carbides is bainite. In addition, in a case in which martensite is
tempered martensite, since tempered martensite includes carbides therein, there is a
possibility that martensite may be misidentified as bainite. However, in steel
according to the present embodiment, it has been clarified that, when the total volume
percentage of martensite, bainite, pearlite, and residual austenite is set to be 5%,
favorable reduction in area can be obtained, and thus the influence of the
misidentification of martensite and bainite having influences on the final form of the
steel according to the present embodiment is extremely small. Meanwhile, the
volume percentage of ferrite is desirably set to equal to or higher than 70%.
[0060]
The volume percentage of residual austenite is measured by X-ray diffraction.
A strained layer on the surface of a sample, which is obtained by finishing the
observation surface to be a mirror-like surface in the above-described order, is
removed using electro-polishing, thereby preparing a sample for measuring residual
austenite. Electro-polishing is carried out using a 5% perchloric acid-acetic acid
solution by applying a voltage of 10 V. Cu is selected as an X-ray tube, and the
volume percentage of residual austenite is obtained on the basis of strengths on
individual planes of (200), (220), and (31 1) of austenite and of (200) and (21 1) of
ferrite.
[0061]
Carbides are observed using a scanning electron microscope. Samples for
structural observation are prepared by finishing observed sections to be mirror-like
surfaces by wet-polishing using Emery paper and polishing using diamond abrasive
grains having a particle size of 1 pm and then canying out etching using a saturated
picric acid alcohol solution. The observation magnification is in a range of 1,000
times to 10,000 times, and, in the present embodiment, 16 visual fields including equal
to or more than 500 carbides are selected on ihe structural observation surface at a
magnification of 3,000 times, and structural images are obtained. From the obtained
structural images, the areas of the respective carbides in this region are measured in
detail using image analysis software represented by Win ROOF manufactured by
Mitani Corporation. The circle-equivalent diameters ("circle-equivalent
diametern=2x("area"/3. 14)'") of the respective carbides are obtained from the areas of
the respective carbides, and thc average value thereof is used as the carbide particle
diameter. Meanwhile, in order to suppress the influence of measurement errors
caused by noise, carbides having an area of equal to or smaller than 0.01 pm2 are
excluded from evaluation subjects.
[0062]
A preferred rangc of the carbide particle diameter is 0.30 pm to 1.50 pm. In
a case in which the carbide particle diameter is smaller than 0.30 pm, the ferrite grain
size becomes too small, and thus the lower limit of the carbide particle diameter is set
to 0.30 pm. When the carbide particle diameter exceeds 1.50 pm, it becomes easy for
voids to be initiated in the vicinities of carbides during the distortion of the steel sheet,
and deformability is degraded, and thus the upper limit of the carbide particle diameter
is set to 1.50 pm. In addition, carbides having a ratio of the long-axis length to the
short-axis length of equal to or greater than 3 are determined as needle-like carbides,
and carbides having a ratio of the long-axis length to the short-axis length of smaller
than 3 are determined as spherical carbides. The value obtained by dividing the
number of spherical carbides by the number of all carbides is used as the spheroidizing
ratio of carbides (cemenlite and the like).
[0063]
The presence or absence of crystal interfaces at which an orientation
difference is equal to or greater than 5" is investigated using EBSD. Samples for
evaluation are cut out using a discharging wire processing machine from a place, to
which strain is not imparted, of a steel strip and a cut sheet cut out from a steel strip or
a blank sheet obtained from a steel strip by punching, and planes thereof perpendicular
to the surface of the steel sheet are used as observed sections. Since the measurement
accuracy of EBSD is affected by the flatness of the observation surface and strain
imparted by polishing, the observation surface is finished to be a mirror-like surface by
wet-polishing and diamond abrasive grain polishing, and thcn polishing for removing
strain is carried out on the observation surface. Strain-removing polishing is carried
out using an oscillatory polishing device (VibroMet 2 manufactured by Buhler AG)
under conditions of an output of 40% and a polishing time of 60 min. When SEMEBSD
is used, the device type of SEM and Kikuchi-line detector are not particularly
limited. In a quarter thickness layer, four visual fields are measured at measurement
step intervals of 0.2 pm in a region 100 in the sheet thickness direction and 100 pm
in the sheet width direction, and an orientation difference regarding crystal interfaces
present in individual cementite are measured and the number of particles having a
crystal interface of equal to or greater than 5" are counted from the obtained map
information of crystal orientations. Measurement data are preferably analyzed using
OIM analysis software manufactured by TSL, and, in order to eliminate the influence
on data of measurement errors caused by noise, cleanup is not carried out, and data
having a coincidence index (CI value) of equal to or lower than 0.1 is excluded in the
analysis.
COO641
When the ferrite grain size in the structure after cold-rolled-sheet-annealing is
5 pm to 60 pm, it is possible to suppress reduction in area being decreased during
distortion at a high strain rate. When the ferrite grain size is smaller than 5 pm,
deformability is degraded, and thus the lower limit of the ferrite grain size is set to 5
pm. In addition, when the ferrite grain size exceeds 60 pm, satin is generated on the
surface in the initial phase of distortion, and breakage is accelerated by surface
irregularity formed thereon as starting point, thereby decreasing reduction in area.
Therefore, the upper Illnit of the ferrite grain size is set to equal to 01 smaller than 60
pm. The ferrite grain sizc is measured by finishing the observation surface to be a
mirror-like surface by polishing in the abovc-described order, etching the surface using
a 3% nitric acid-alcohol solution, observing the structure using an optical microscope
or a scanning electron microscope, and applying a line segment method on a captured
image. The ferrite grain size is preferably 10 pm to 50 pm.
[0065]
Subsequently, a method for measuring reduction in area during distortion at a
high strain rate will be described.
[0066]
In order to distort the steel sheet at a strain rate of 10 mmlsec and measure
reduction in area during breakage, it is necessary to use a special test specimen having
1.5 mm-long parallel portions which are shown in FIG. 1. When a tensile test is
carried out on the special test specimen having the 1.5 mm-long parallel portions at a
stroke rate of 900 mmlminute, it becomes possible to impart strain to the parallel
portions in the test specimen at a strain rate extremely close to 10 mmlsec. In
addition, in order to accurately evaluate the behaviors of the fracture of the steel sheet
which may occur during shaping into actual components, it is necessary to strictly
manage the ratio of the thickness to the width of the parallel portions in the tensile test
specimen. During the drawing distortion of the tensile test specimen, neckiug
distortion occurs in two directions of the thickness direction and the width direction.
It is needless to say that, when breakage occurs during the shaping of actual
components, necking distortion in the thickness direction is a dominant factor of the
breakage, and the influence of necking distortion in the width direction is extremely
small. Therefore, in evaluation in which a tensile test specimen is used, it is
necessary to remove the influence of necking distortion in the width direction, and thus
it is ncccssary to set the ratio of the width of the parallel portion to the thickness of the
parallel portion to bc equal to or greater than 2. The ratio of width to thickness is
preferably greater, more preferably equal to or greater than 4, and still more preferably
equal to or greater than 6. In addition, reduction in area is calculated from a change
in the thickness before and afler tensile breakage using Equation (1)
"Reduction in area (%)"=(("the sheet thickness before the testn-"the sheet
thickness after breakagem)/"the sheet thickness before the test")^ 100 *a- (1)
Meanwhile, the thickness before the test is obtained by measuring the
thickness at the central portion in the width direction of the parallel portion and the
thicknesses at two points respectively 1 mm away from the central portion in a
direction that is perpendicular to the longitudinal direction and is parallel to the width
direction using a micrometer and averaging the measurement values at the three points.
The thickness of the sample afler breakage is measured using, for example, a
microscope (VHX-1000) manufactured by Keyence Corporation. Similar to the
measurement before the test, the thicknesses at the width central portions and the
thicknesses at locations 1 mm away from the central portions in the width direction in
each of the broken surfaces of the sample that has been divided into two pieccs due to
breakage are respectively measured, and the average of the measurement values at six
points is used as the thicltness after the test. Samples exhibiting high reduction in
area of equal to or greater than 10% in the above-described test were evaluated as
samples exhibiting "excellent reduction in area".
[0067]
Next, a method for manufacturing a stecl sheet according to the present
embodiment will be described.
[0068]
The technical concept of the method for manufacturing a steel sheet according
to the present embodiment is to collectively manage the conditions of hot-rolling and
anncaling using a material having the above-described composition ranges.
[0069]
The characteristics of a specific method for manufacturing a steel sheet
according to the present embodiment will be described below.
[0070]
In hot-rolling, when a slab having predetermined composition is is hot-rolled
directly after continuously-casting as per ordinary method or hot-rolled after temporary
cooling and heating, finish hot-rolling is terminated in a temperature range of 600°C to
lower than l,OOO°C. The finishing-rolled steel strip is cooled on a run-out table
(ROT) at a cooling rate of 10 "Clsecond to 100 "C/second and then is coiled in a
temperature range of 350°C or more and less than 700°C, thereby obtaining a hotrolled
coil. Hot-rolled-sheet-annealing is carried out on the hot-rolled coil,
subsequently, cold-rolling is carried out at a cold-rolling reduction ratio of 10% to 80%,
and furthermore, cold-rolled-sheet-annealing is carried out, thereby obtaining a
middlehigh carbon stecl sheet exhibiting excellent reduction in area during distortion
at a high strain rate.
[0071]
Hereinafter, the method for manufacturing a steel sheet according to the
present embodiment will be specifically described.
[0072]
(Hot-rolling)
When a slab (billet) having predetelmined composition is continuously cast, is
heated directly or after temporary cooling, and then is hot-rolled, finish hot-rolling is
completed in a temperature range of 600°C or more and less than 1 ,00O0C, and the
obtained steel strip is coiled in a temperature range of 350°C or more and less than
700°C.
[0073]
The heating temperature of the slab is 950°C to 1,250°C, and the heating time
is set to 0.5 hours to three hours. In a case in which the heating temperature exceeds
1,250°C or the heating time exceeds three hours, decarburization from the slab surface
layer becomes significant, and the hardness of the surface layer decreases even when a
heat treatment of quenching is carried out thereon, and thus wear resistance and the
like necessary for components cannot be obtained. Therefore, the upper limit of the
heating temperature is set to equal to or lower than 1,25OoC, and the upper limit of the
heating time is set to equal to or shorter than three hours. In addition, in a case in
which the heating temperature is lower than 950°C or the heating time is shorter than
0.5 hours, micro segregation or macro segregation formed during casting is not
resolved, and regions in which alloy elements such as Si and Mn locally thicken
remain in steel products, and these regions cause a decrease in reduction in area during
distortion at a high strain rate. Therefore, the lower limit of the heating temperature
is set to equal to or higher than 95OoC, and the lower limit of the heating time is set to
equal to or longer than 0.5 hours.
[0074]
Finish hot-rolling is preferably ended at 600°C to l,OOO°C. When the finish
hot-rolling temperature is lower than 600°C, an increase in the deformation resistance
of steel products significantly increases the rolling load and, furthermore, increases the
roller wear amount, and thus productivity is decreased. Therefore, the finish hotrolling
temperature is set to equal to or higher than 600°C. Ln addition, when the
finish hot-rolling temperature exceeds 1,00O0C, thick scales are generated on the steel
sheet during the passing ofthe steel sheet through a run-out table, the scales serve as
oxygen sources, and grain boundaries of ferrite or pearlite are oxidized after coiling,
thereby forming fine protrusions and recesses on the surface. Since the steel sheet
breaks from the fine protrusions and recesses as starting points in an early phase during
distortion at a high strain rate, the fine protrusions and recesses cause a decrease in
reduction in area. Furthermore, when the finish hot-rolling temperature exceeds
1,00O0C, segregation of alloy elements such as Si and Mn into austenite grain
boundaries after the finish hot-rolling is accelerated, and the concentrations ofthe alloy
elements in austenite grains decrease, and thus carbides agglomerate during hot-rolledsheet-
annealing and cold-rolled-sheet-annealing at portions at which the concentrations
of the alloy elements are low, and the proportion of the number of carbides having a
crystal interface increases. Therefore, the finish hot-rolling temperature is set to
equal to or lower than 1 ,00O0C.
100751
The cooling rate of the steel strip on ROT after the finish hot-rolling is set to
10 "Clsecond to 100 OCIsecond. In a case in which the cooling rate is slower than
10 "Clsecond, the cooling rate is slow, and thus the growth of ferrite is accelerated, and
a structure in which ferrite, pearlile, and bainite are laminated in the sheet thickness
direction of the steel strip is formed in the hot-rolled sheet. The above-described
structure remains even after cold-rolling and annealing and causes a decrease in the
reduction in area ofthe steel sheet, and thus the cooling rate is set to equal to or faster
than 10 "C/secoud. In addition, when the steel strip is cooled at a cooling rate
cxcecding 100 "Clsecond throughout the entire sheet thickness, the outermost surface
part is excessively cooled, and low temperature transformation structures such as
bainite and martensite are generated. When a coil cooled to a range of 100°C to room
temperature after coiling is discharged, fine cracks occur at thc above-described low
temperature transfor~nations tructures. In the subsequent pickling and cold-rolling, it
is difficult to remove the cracks, and the cracks decrease the reduction in area of the
steel sheet that has been subjected to the cold-rolled-sheet-annealing. Therefore, the
cooling rate is set to equal to or slower than 100 "Cisecond. Meanwhile. the cooling
rate determined above refers to the cooling power received from cooling facilities
between individual water injection zones from a timing when a steel strip that has been
subjected to finish hot-rolling is water-cooled in a water injection zone after passing
through a non-water injection zone and a timing when the steel strip is cooled on ROT
to the target coiling temperature, and does not refer to the average cooling rate applied
hom the start of water injection to coiling in which a coiling device is used.
COO761
The coiling temperature is set to 350°C to 700°C. When the coiling
temperature is lower than 35OoC, austenite which has remained untransformed during
the finishing rolling transforms to martensite, fine ferrite and cementite are maintained
even after the cold-rolled-sheet-annealing, and reduction in area is decreased, and thus
4 the coiling temperature is set to be equal to or higher than 350°C. In addition, when
I:
the coiling temperature exceeds 700°C, untransformed austenite transforms to pearlite
i having coarse lamellar, and bulky needle-like cementite remains even after the cold-
$I rolled-sheet-annealing, and thus reduction in area is decreased. Therefore, the coiling
i
I temperature is set to be equal to or lower than 700°C.
Box-annealing is carried out on the hot-rolled coil manufactured under the
above-described conditions directly or after pickling. The annealing temperature is
set to 670°C to 770°C, and the retention time is set to one hour to 100 hours.
[0078]
The box-annealing temperature is preferably set to 670°C to 770°C. When
the annealing temperature is lower than 670°C, ferrite grains and carbide particles do
not sufficiently coarsen, and reduction in area is decreased during distortion at a high
strain rate. Therefore, the annealing temperature is set to be equal to or higher than
670°C. In addition, when the annealing temperature exceeds 770°C, the structural
ratio of ferrite during the annealing in two-phase region of ferrite and austenite is
excessively small, and thus it is not possible to avoid the generation of pearlite having
a large lamellar spacing even when the steel sheet is cooled to room temperature at an
extremely slow cooling rate of 1 'Clhr during the box-annealing, and the spheroidizing
ratio after the cold-rolled-sheet-annealing is decreased, and thus reduction in area
during distortion at a high strain rate is decreased. Therefore, the annealing
temperature is set to be equal to or lower than 770°C. The annealing temperature is
preferably 685°C to 760°C.
[0079]
The retention time of the box-annealing is preferably set to one hour to 100
hours. When the retention time is shorter than one hour, carbides do not sufficiently
spheroidize during the hot-rolled-sheet-annealing, and the spheroidizing ratio is low
cven after the cold-rolled-sheet-annealing, and thus reduction in area is decreased.
Therefore, the retention time of box-annealing is set to equal to or longer than one hour.
Under a condition in which the retention time exceeds 100 hours, productivity
degrades, and interfaces are formed due to carbides being combined together or
coming into contact with each other, and thus the retcntion time of box-annealing is set
to equal to or shorter than 100 hours. The lower limit of the retention time of boxannealing
is preferably two hours and more preferably five hours, and the upper limit
thereof is preferably 70 hours and more preferably 38 hours.
[OOSO]
Meanwhile, the atmosphere for the box-annealing is not particularly limited
and may be any one of an atmosphere of equal to or higher than 95% of nitrogen, an
atmosphere of equal to or higher than 95% of hydrogen, and the atmospheric
atmosphere.
[008 11
Next, the reasons for carrying out the cold-rolling at a cold-rolling reduction
of 10% to 80% will be described. In the above-described hot-rolling and hot-rolledsheet-
annealing, the coil after hot-rolled-sheet-annealing, which has been subjected to
picMing before or after the hot-rolled-sheet-annealing, is cold-rolled at a cold-rolling
reduction of 10% to 80%. Ln a case in which the cold-rolling reduction is lower than
lo%, during the cold-rolled-sheet-annealing, the number of nuclei for the
recrystallization of ferrite is small, the ferrite grain size coarsens, and the steel sheet
breaks from satin generated on the steel sheet surface during distortion at a high strain
rate as starting points, and thus reduction in area is decreased. Therefore, the lower
limit of the cold-rolling reduction is set to 10%. Tn addition, when the cold-rolling
reduction exceeds SO%, the number of nuclei for the recrystallization of ferrite is large,
the grain size of ferrite obtained after the cold-rolled-sheet-annealing becomes too
small, and deformability degrades, and thus reduction in area is decreased during
distortion at a high strain rate. Therefore, the upper limit of the cold-rolling reduction
is set to 80%.
[0082]
When the cold-rolled-sheet-annealing is carried out on a steel strip that has
been cold-rolled at the above-described cold-rolling reduction, it is possible to obtain a
middlethigh carbon steel sheet exhibiting excellent reduction in area during distortion
at a high strain rate.
[0083]
Meanwhile, during the cold-rolled-sheet-annealing, the diffusion frequency of
individual elements in steel increases due to the presence of lattice defects such as
dislocation introduced by the cold-rolling. Therefore, during the cold-rolled-sheetannealing,
a change in which carbide particles do Ostwald growth, coarsened carbide
particles come into contact with each other and thus form a single particle, and crystal
interfaces are formed in the carbide particle, is likely to occur. Long-time annealing
allows the above-described change of carbide particles to be more significant, and thus
the cold-rolled-sheet-annealing is desirably carried out in a continuous annealing
furnace.
[0084]
Subsequently, the conditions for the cold-rolled-sheet-annealing by
continuous annealing will be described. The continuous annealing is desirably
carried out at an annealing temperature of 650°C to 780°C for a retention time of 30
seconds to 1,800 seconds. When the annealing temperature is lower than 65OoC, the
size of ferrite obtained after the cold-rolled-sheet-annealing is small, and deformability
is low, and thus reduction in area during distortion at a high strain rate is decreased.
Therefore, the lower limit of the annealing temperature is set to 650°C. In addition,
when the annealing temperature exceeds 780°C, the ratio of austcnite being generated
during the annealing excessively increases, and thus it is not possible to wppress the
generation of martcnsite, bainite, pearlite, and residual austenite after the cooling, and
reduction in area is decreased. Therefore, the upper limit of the annealing
temperature is set to 780°C. Furthermore, when the retention time is shorter than 30
seconds, the size of ferrite obtained after the cold-rolled-sheet-annealing becomes
small, and thus reduction in area is decreased. Therefore, the lower limit ofthe
retention time is set to 30 seconds. In addition, when the retention time exceeds
1,800 seconds, in a process in which carbide particles grow during the cold-rolledsheet-
annealing, carbide particles come into contact with each other, crystal interfaces
are formed in the particles, and reduction in area is decreased. Therefore, the upper
limit of the annealing time is set to equal to or shorter than 1,800 seconds.
Meanwhile, the heating rate, the cooling rate, and the temperature of an OA zone
(over-ageing zone) during the cold-rolled-sheet-annealing are not particularly limited,
and, in studies of tests according to the present embodiment, it is confirmed that, under
conditions of a heating rate of 3.5 'Clsecond to 35 "Clsecond, a cooling rate of
1 "Clsecond to 30 OCIsecond, and the temperature of the OA zone of 250°C to 450°C,
intended forms of the steel sheet according to the present embodiment are sufficiently
obtained.
[0085]
According to the above-described method for manufacturing a steel sheet of
the present embodiment, it is possible to obtain a middlehigh carbon steel sheet
exhibiting excellent formability when deformation processing such as deep drawing,
hole expanding, thickening, or thinning or cold forging in which the above-described
processes are combined together is carried out at a high strain rate by providing a
structure including ferrite and carbides as main bodies, setting the total volume
percentage of martensite, bainite, pearlite, and rcsidual austenite to be equal to or lower
than 5.0%, setting the spheroidizing ratio of carbide particles to be 70% to 99%, and
setting the proportion of the number of the carbide particles including a crystal
interface at which an orientation difference is equal to or greater than 5" in the carbide
particles to be equal to or lower than 20% of the total number of carbide particles.
EXAMPLES
[0086]
Next, the effects of the present invention will be described using examples.
[0087]
The levels of the examples are examples of conditions for carrying out the
present invention which were employed to confirm the feasibility and effects of the
present invention, and the present invention is not limited to these condition examples.
The present invention allows employment of a variety of conditions within the scope
of the gist of the present invention as long as the object of the present invention is
achieved.
[008S]
Continuous cast pieces (steel ingots) having a composition shown in Table 1
were heated at 1,140°C for 1.6 hours and then were hot-rolled, thereby obtail~ing2 50
mm-thiclc slabs. The slabs were roughly hot-rolled to a thickness of 40 mm, rough
bars, which are materials for finishing hot-rolling, were heated by 36OC so as to initiate
finish hot-rolling, the rough bars were finishing-hot-rolled at 880°C, then, were cooled
to 520°C on ROT at a cooling rate of 45 "Clsecond, and were coiled at 5 10°C, thereby
manufacturing hot-rolled coils having a sheet thickness of 4.6 mm. The hot-rolled
coils were pickled and were loaded into a box-type annealing furnace, the atmosphere
was controlled to be 95% hydrogen-5% nitrogen, the hot-rolled coils were heated from
room temperature to 500°C at a heating rate of 100 "Clhour, and were held at 500°C
for three hours, thereby evening the temperature distributions in the coils. After that,
the hot-rolled coils were heated to 705°C at a heating rate of 30 "Chour, were further
held at 705°C for 24 hours, and then were cooled to room temperature in the furnace.
The coils which had been subjected to hot-rolled-sheet-annealing were cold-rolled at a
rolling reduction of 50% and cold-rolled-sheet-annealing in which the coil were held at
720°C for 900 seconds was carried out, and temper rolling was carried out at a rolling
reduction of 1.2%, thereby producing samples for characteristic evaluation. The
structure and the reduction in area during distortion at a high strain rate of the samples
were measured using the above-described methods.
[0089]
Tables 2-1 and 2-2 show the evaluation results of the reduction in area during
distortion at a high strain rate of the manufactured samples. As shown in Tables 2-1
and 2-2, in all of Invention Examples No. B-1, C-1, D-1, E-1, F-1, G-1, H-l, I-1, J-1,
M-1, N-1, P-1, Q- 1, R-1, S-1, U-1, X-1, Y-1, Z-1, AA-1, AB-I, and AC-1, the total
volume percentage of martensite, bainite, pearlite, and residual austenite was equal to
or lower than 5%, the spheroidizing ratio of carbide particles was equal to 70% to 99%,
and the proportion of thc number of carbide particles including a crystal intet lace at
which an orientation difference is equal to or greater than 5" in carbide particles was
equal to or lower than 20% of the total number of the carbide particles, and excellent
reduction in area during distortion at a high strain rate was exhibited.
[0090]
In contrast, in Comparative Example A-1, the proportion of carbides having
crystal interfaces was low and excellent reduction in area during distortion at a high
strain rate was exhibited, hut the C content was low, and high-strengthening was not
possible in the quenching step for producing prodncts, and thus the steel sheet was
I
I evaluated as fail. In Comparative Example I<-1, the Mn content was low, the Oswald
growth of carbides was accelerated during the cold-rolled-sheet-annealing, and the
proportion of carbides having crystal interfaces increased, and thus the reduction in
area was decreased. In Comparative Example L- 1, the P content was high, ferrite
grain boundaries embrittled, and fissures were initiated and propagated from ferrite
grain boundaries during distortion at a high strain rate, and thus the reduction in area
was decreased. In Comparative Example 0-1, the Mn content was high,
spheroidization of carbides during the hot-rolled-sheet-annealing and the cold-rolledsheet-
annealing was suppressed, and fissures were initiated and propagated from
needle-like carbides during distortion at a high strain rate, and thus the reduction in
area was decreased. In Comparative Example T-1, the Si content was low, the
Oswald growth of carbides was accelerated during the cold-rolled-sheet-annealing, and
the proportion of carbides having crystal interfaces increased, and thus the reduction in
area was decreased. In Comparative Example V-1, the S content was high, a number
of coarse inclusions such as MnS were present in steel, and fissures were initiated and
propagated from the inclusions as starting points, and thus the reduction in area was
decreased. In Comparative Example W-1, the Si content was high, it became difficult
for austenite generated during the cold-rolled-sheet-annealing to do ferritic
transfornation during cooling, and bainitic and pearlitic transformation was promoted,
and thus the structural proportion of those other than ferrite and carbides increased,
whereby stress accumulated in ferrite grain boundaries, and the reduction in area was
decreased. In Comparative Example AD-1, the C content and the volume percentage
of carbides were high, it was not possible to control the proportion of the number of
carbides having crystal interfaces to be equal to or lower than 20%, and the reduction
in area was decreased.
[Table 11
[Table I]
* UNDEllLlNED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
[Table 2-11

LO0941
Next, in order to investigate the ranges of the permissible contents of other
elements, continuous cast pieces (ingots) having a composition shown in Tables 3-1,3-
2, 3-3,4-1,4-2, and 4-3 were heated at 1,180°C for 0.7 hours and then were hot-rolled,
thereby obtaining 250 mm-thick slabs. The slabs were roughly hot-rolled to a
thiclcness of 45 mm, rough bars, which is materials for finishing hot-rolling, were
heated by 48°C so as to initiate finish hot-rolling, the rough bars were finishing-hotrolled
at 870°C, then, were cooled to 5 10°C on ROT at a cooling rate of 45 "Cisccond,
and were coiled at 500°C, thereby manufacturing hot-rolled coils having a sheet
thickness of 2.6 mm. The hot-rolled coils were pickled and were loaded into a boxtype
annealing furnace, the atmosphere was controlled to be 95% hydrogen-5%
nitrogen, the hot-rolled coils were heated from room temperature to 500°C at a heating
rate of 100 "Cihour, and were held at 500°C for three hours, thereby evening the
temperature distributions in the coils. After that, the hot-rolled coils were heated to
705°C at a heating rate of 30 "Cihour, were further held at 705'C for 24 hours, and
then were cooled to room temperature in the furnace. The coils which had been
subjected to hot-rolled-sheet-annealing were cold-rolled at a rolling reduction of 50%
and cold-rolled-sheet-annealing in which the coils were held at 700°C for 900 seconds
was carried ouf and temper rolling was carried out at a rolling reduction of 1.0%,
thereby producing samples for characteristic evaluation.
[0095]
Tables 5-1 and 5-6 show the evaluation results of the reduction in area during
distortion at a high strain rate of the manufactured samples. As shown in Tables 5-1
and 5-6, in all of Invention Examples No. AE-1, AF-1, AL-I, AM-1, AN-1, All-1, AS-
1, AV-1, AW-1, AX-1, BC-1, BD-1, BF-1, BH-1, BI-1, BJ-I, BK-1, BM-I, BN-I, and
BT- 1, the total volume percentages of martensite, bainite, pearlite, and residual
austenite were equal to or lower than 5% (including 0.0%), the spheroidizing ratios of
carbide particles were 70% to 99%, and the proportions of the number of carbide
particles including crystal interface at which an orientation difference is equal to or
greater than 5" in carbide particles were equal to or lower than 20% of the total number
of the carbide particles, and excellent reductions in area during distortion at a high
strain rate were exhibited.
[0096]
In contrast, in Comparative Examples AG-1, AH-1, AO-1, AT-1, AU-I, AZ-1,
BA-1, BB-1, BO-1, and BS-1, the contents of Ce, Ca, Y, Al, Mg, As, Zr, Sn, Sb, and
La were respectively high, and thus grain boundaries of ferrite embrittled, and the
reductions in area were decreased during distortion at a high strain rate. In
Comparative Examples AI-I, AJ-1, AK-I, AQ-I, BE-1, BG-1, BL-I, BQ-I, and BR-I,
the contents of Nb, W, Ti, Ni, Cr, Mo, V, Cu, and Ta were high, spheroidization of
carbides during the hot-rolled-sheet-annealing and the cold-rolled-sheet-annealing was
suppressed, and fissures were initiated and propagated from needle-like carbides
during distortion at a high strain rate, and thus the reduction in area was decreased. In
Comparative Example AP-1, the N content was high, it became difficult for austenite
generated during the cold-rolled-sheet-annealing to do ferritic transformation during
cooling, and bainitic and pearlitic transformation was promoted, and thus the structural
proportion of those other than femte and carbides increased, whereby stress
accumulated in ferrite grain boundaries, and the reduction in area was decreased. In
Comparative Example AY-1, the 0 content was high, coarse oxides were formed in
steel, and fissures were initiated and propagated from the coarse oxides as starling
points during distortion at a high strain rate, and thus the reduction in area was
decreased. In Comparative Example BP-1, the B content was high, and coarse Fe-Bcarbides
were generated in steel, and thus fissures were initiated and propagated from
the Fe-B-carbides as starting points, and thus the reduction in area was decreased.
[0097]
[Table 3 - 1 ]
[0098]
[Table 3-21
r-.
0
-0ww
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
[Table 3-31
- 2
%
Y W -
P w
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
- 7
0
e ,..A
L...
I * UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.

[Table 4-31
INVENTION EXAMPLE
* UNDERLINED BOLD NTJMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
[0103]
[Table 5-11
[Table 5-1 1
AP-I 79.2 0.38 8.7 0.0% 4.8% 0.8% 0.3% 5.9%
AQ-l 82.9 0.46 9.6 0.0% 0.0% 0.0% 0.0% 0.0%
AR-I 87.6 0.48 9.8 0.0% 0.0% 0.0% 0.0% 0.0%
* UNDEKINED BOLD NIJMERICAL VALUES INDICATE EXAMPLES OUTSIDE
THE INVENTION RANGES.
[0 1041
[Table 5-21
AE- 1
AF-1
AG-1
AI-1 I 2.1 I 6.9 I COMPARATIVEEXAMPLE
AH- 1
1.3
2.2
4.8
AN-I I 2.7 1 12.5 1 INVENTIONEXAMPLE
2.0
AJ-I
AO-1 I 4.6 1 5.7 1 COMPARATIVE EXAMPLE
12.6
12.6
5.3
INVENTION EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
-6.9
3.3
COMPARATIVE EXAMPLE
AP-1
* UNDERLINED BOLD NUMERICAL. VALUES TNDICATE EXAMPLES
OUTSIDE THE INVENTION RANGES.
-6.2
AQ-l
[Table 5-31
COMPARATIVE EXAMPLE
4.5 -4.9 COMPARATIVE EXAMPLE
3.2
AR-1
-6.4 COMPARATIVE EXAMPLE
3.0 12.5 INVENTION EXAMPLE
[Table 5-31
STKUCTURAL PROPORTION OF THOSE
OTHER THAN FERRITE AND CARBlDES
* UNDERLINE11 BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE
THE INVENTION RANGES.
AW-I
AX-1
AY-1
[0106]
[Table 5-41
86
92.1
81.9
0.46
0.48
0.45
I
0.0%
0.0%
0.0%
9.1
9.2
8.8
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
NOTE
3.2 1 12.5 1 INTENTION EXAMPLE
5.4 1 2.5 1 COMPARATIVEEXAMPLE
3.4 1 ZZ 1 COMPARATIVEEXAMPLE
AV-1
AW-1
AX- 1
AY-1
AZ- 1
BA-1
BA-I
* UNIIERLWED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE THE INVENTION RANGES.
BC-1
BD-1
BE-1
BF-1
[Table 5-51
3.9
3.7
3.5
4.1
5.2
4.7
7.1
4.7
4.4
13.8
4.4
12.5
12.5
12.5
-3.1
-5.2
-8.3
4.2
INVENTION EXAMPLE
INVENTION EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
COMPARATIVE EXAMPLE
COMPARATIVE EXAMPLE
COMPARATIVE EXAMPLE
12.5
12.5
-4.2
12.5
INVENTION EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
INVENTION EXAMPLE
STRUCTURAL PROPOR1 ION 01. THOSt
OTILER TIlAN FERRITE AND CARBIDES
W
W W
dt: 4
3 2
d
2
m 8
"2 P
BM-I
BNN-I
BO-1
THE INVENTION RANGES.
[OlOS]
[Table 5-61
BS-l 93.1 0.46 5.8 I 0.0% I 0.0% I 0.0% I 0.0% 0.0%
85.9
80.4
86.1
0.5%
0.46
0.34
0.42
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE
0.0%
7.4
5.0
6.5
0.0%
0.0%
0.0%
0.0%
0.1% 0.6%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
0.0%
[Table 5-61
NOTE
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE THE INVENTION RANGES
[0 1091
Subsequently, in order to investigate the influences of the manufacturing
conditions, slabs having compositions No. B, C, D, E, F, G, H, I, J, M, N, P, Q, R, S, U,
X, Y, Z, AA, AB, AC, AE, AF, AL, AM, AN, AR, AS, AV, AW, AX, BC, BD, BF,
BH, BI, BJ, BK, BM, BN, and BT shown in Tables l,3-1 to 3-3 and 4-1 to 4-3 were
cast and temporarily cooled, and then hot-rolled steel strips having a sheet thickness of
3.5 mm were manufactured under the slab heating conditions and the hot-rolling
conditions shown in Tables 6-1-1,6-1-2,6-2-1,6-2-2, 7-1-1,7-1-2,7-2-1,7-2-2, 8-1-1
to 8-1-3, 8-2-1 to 8-2-3, 9-1-1 to 9-1-3, and 9-2-1 to 9-2-3 (hereinafter, simply denoted
as Tables 6, 7, 8, and 9), and hot-rollcd steel annealing, pickling, cold-rolling, and
cold-rolled-sheet-annealing were carried out, thereby manufacturing samples for
characteristic evaluation.
[Ol lo]
Tables 6,7,8, and 9 also show the evaluation results of the reduction in area
during distortion at a high strain rate of the manufactured samples. As shown in
Table 8, in all of Invention Examples No. B-2, C-2, D-2, E-2,J-2, N-2, Q-2, X-2, Y-2,
2-2, AB-2, AC-2, AL- 2, AN-2, AS-2, AV-2, BC-2, BD-2, BH-2, BI-2, BJ-2, BN-2, F-
3, G-3, H-3,I-3, M-3, N-3, P-3, R-3, S-3, U-3, AA-3, AB-3, AE-3, AF-3, AM-3, AR-
3, AW-3, AX-3, BF-3, BK-3, BM-3, and BT-3, the total volume percentages of
martensite, hainite, pearlite, and residual austenite were equal to or lower than 5%, the
spheroidizing ratios of carbide particles were 70% to 99%, and the proportions of the
number of carbide particles including crystal interfaces at which an orientation
difference is equal to or greater than 5" in carbide particles were equal to or lower than
20% of the total number of the carbide particles, and excellent reductions in area
during distortion at a high strain rate were exhibited.
[Olll]
In contrast, in Comparative Examples AA-2, BK-2, C-3, and BJ-3, as shown
in Tables 6 and 7, the finish hot-rolling temperatures were high, the proportions of the
number of carbides having crystal interfaces increased, bulky scales generated from the
coiling to the cooling served as oxygen supply sources, grain boundaries were oxidized
after the coiling, and fine cracks were generated on the surface, whereby fissures
propagated from cracks in the surface layer as starting points during distortion at a high
strain rate, and thus the reductions in area were decreased. In Comparative Examples
R-2, BM-2, X-3, and BC-3, the finish hot-rolling temperatures were low, and, when
rolling was carried out by involving scales during the hot-rolling, protrusions and
recesses were formed on the surfaces of the steel sheets, and fissures were initiated and
propagated from the protrusions and recesses on the surface layer as starting points
during distortion at a high strain rate, and thus the reductions in area were decreased.
In Comparative Examples U-2, AR-2, Y-3, and AL-3, the coiling temperatures were
high, needle-like carbides having a large thickness were generated in the hot-rolled
sheets, and spheroidization of the needle-like carbides did not proceed even after the
cold-rolled-sheet-annealing, and thus fissures were initiated and propagated from the
needle-like carbides as starting points, and thus the reductions in area were decreased.
In Comparative Examples H-2, AM-2, 4-3, and BI-3, the coiling temperatures were
low, the structures of the hot-rolled sheet were fine, and the structures after the coldrolled-
sheet-annealing were also fine, and thus deformability degraded, and the
reductions in area were decreased during distortion at a high strain rate.
[OllZ]
In Comparative Examples G-2, AE-2,J-3, and BD-3, as shown in Tables 6
and 7, the cold-rolling reductions were high, and thus the structures after the coldrolled-
sheet-annealing hccame fine, deformability degraded, and the reductions in area
were decreased. In Comparative Examples S-2, AW-2, AC-3, and BH-3, the coldrolling
reductions were low, and thus the ferrite grain sizes after the cold-rolled-shectannealing
became coarse, satin was generated on the surface layer during distortion at
a high strain rate, and fissures were initiated and propagated from the protrusions and
recesses formed on the surface, and the reductions in area were decreased. In
Comparative Examples M-2, BT-2,Z-3, and AS-3, the temperatures of the cold-rolledsheet-
annealing were high, and thus the phase ratios of austenite generated during
annealing became high, and it was not possible to suppress martensitic, bainitic, and
pcarlitic transformation in the cooling process, and thus the reductions in area were
decreased during distortion at a high strain rate. In Comparative Examples P-2, BF-2,
E-3, and BN-3, the temperatures of the cold-rolled-sheet-annealing were low, and
ferrite grain boundaries were fine, and thus deformability degraded, and the reductions
in area were decreased during distortion at a high strain rate. In Comparative
Examples 1-2, AX-2, D-3, and AN-3, the cold-rolled-sheet-annealing times were long,
carbides particles came into contact with each other in the coarsening proccss, and
crystal interfaces were formed in the particles, and thus the reductions in area were
decreased. In Comparative Examples F-2, AF-2, B-3, and AV-3, the cold-rolledsheet-
annealing times were short, and ferrite was fine, and thus deformability degraded,
and the reductions in area were decreased during distortion at a high strain rate.
[0113]
[Table 6-1-11
:Table 6-1-11
HOT ROLLING CONDITIONS
I I I I
W
0 TIMING OF PICKLING
W
€-
I I I I I I
B-2 1 1012 1 2.1 1 727 32 470 BEFORE HOT-ROLLED SHEET ANNEALING
1-2 1 1178 1 1.3 1 850 89 359 BEFORE HOT-ROLLED SHEET ANNEALING
5-2 1 997 1 2.8 1 692 15 353 I BEFORE HOT-ROLLED SHEET ANNEALING
[0115]
[Table 6-2-11
[Table 15-2-11
i HOT ROLLING CONDITIONS
W
TIMING OF PICKLING
W
I- I- f-
I - / 903 1 33 / 432 1 BEFORE HOT-ROLLED SHEET ANNEALING
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES

[Table 7-1-11
HOT ROLLING CONDITIONS - - 1 1 E4 s CD u
4 ..-
,j u
No
E-3
F-3
G-3
H-3
B-3
C-3
D-3
1-3
J-3
M-3
HEATING
TEMPERATURE
("C)
1120
1164
1156
987
I I I
1112
1126
963
1214
1059
N-3
P-3
Q-3
- R-3
HEATING
TIME
0.9
1.0
3.0
1.5
S-3
U-3
X-3
Y-3
0.9
2.0
0.6
1.8
1.1
671
1148
1164
1110
2-3 1 1239
FINISHING
TEMPERATURE
("C)
607
782
840
759
1030
1165
1060
1039
793
-1026
900
682
629
847
32
1.0
0.8
1.5
* UNDERLINED BOLD NUMENCAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
2.1
ROT
RATE
?CIS)
13
98
70
52
2.0
1.6
2.9
1.4
14
79
53
25
42
69
659
630
669
920
809
COILING
TEMPERATURE
("C)
361
437
549
645
AFTER HOT-ROLLED SHEET ANNEALING
876
963
-577
798
TIMING OF PICKLING
445
465
370
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
BEFORE HOT-ROLLED SHEET ANNEALING
356
551
598
99
16
79
57
BEFORE HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
16
20
56
79
650 I AFTER HOT-ROLLED SHEET ANNEALING
429
337
410
BEFORE HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
BEFORE HOT-ROLLED SHEET ANNEALING
381
352
423
-718
BEFORE HOT-ROLLED SHEET ANNEALING
BEFORE HOT-ROLLED SHEET ANNEALMG
AFTER HOT-ROLLED SHEET ANNEALING
AFTER HOT-ROLLED SHEET ANNEALING
No
C-3 1 109.1 1 459.7 1 8.9 1 48.5 1 693 I 99.0 I 34.1 I COMPARATIVE EXAMPLE
/ (OCflu) I PC) I (hr) / (OCW I ('C) I Oir 1
HEATING
RATE
UNTIL T1
B-3 1 90.6 1 537.2 1 9.2 1 78.2 1 711
D-3 1 149.3 1 533.6 1 5.9 1 32.3 1 690
G-3 1 100.6 1 459.3 1 2.7 I 58.7 I 755 I 70.1 I 13.8 I INVENTION EXAMPLE
lST COIL
SO4KING
TEMP. T1
97.2
E-3
44.0 I COMPARATIVE EXAMPLE
18.3
RETENTIO
N TIME AT
T1
56.5 I COMPARATIVE EXAMPLE
31.9
I I
F-3 1 45.0 1 523.7 1 5.2 1 49.4 1 715
1-3
5-3
M-3
HEATING
RATE
UNTIL T2
528.7
H-3 INVENTION EXAMPLE
N-3 1 53.1 1 486.2 1 4.0
29.6
117.4
102.8
142.9
~~~~~ ~ ~
P-3
2NDC OIL
SOAKING
TEMP. T2
5.1
74.3
8.6 749
.
79.9
473.0
477.2
526.6
0-3 1 39.7 1 499.7 1 3.9 1 32.0 1 715 I 92.7 I 40.0 I COMPARATIVE EXAMPLE
91.1
S-3
U-3
X-3
RETENTION
AT T2
59.6
INVENTION EXAMPLE
476.2
71.2
2-3 1 127.1 1 451.1 1 3.7 1 49.3 1 729
7.6
8.7
4.8
501.1
R-3 65.5
124.2
114.4
85.2
.,"L,L,
ROLLING
REDUCTION
740
8.3
45.6
70.1 INVENTION EXAMPLE
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES.
68.8
NOTE
12.5
63.6
6.8
INVENTION EXAMPLE
7.0
485.6 61.4
505.8
492.8
516.3
47.8
12.6
4.5 1 46.5 1 75 1
23.5
702
728
765
14.9
3.7
4.1
7.7
COMPARATIVE EXAMPLE
39.1
715
COMPARATIVE EXAMPLE
47.5
41.2
31.1
706
41.8
50.0
45.1
59.0 64.8
54.6
-88.7
55.2
41.9
747
699
754
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
INVENTION EXAMPLE
56.5
53.8
28.5
81.7
INVENTION EXAMPLE
25.4
51.2
61.4
INVENTION EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
,- ---. - -,
1 HOT ROLLING CONDITIONS
,--. ,--. HEATING HEATING FINISHING
ROT
No COILING 2 2 TEMPERATURE TIME TEMPERATURE
RATE TEMPERATURE
TIMING OF PICKLING
N F s ("C) (hr) ("C) ~ C / S ) ("C)
7
Y AA-3 1183 0.5 805 85 654 AFTER HOT-ROLLED SHEET ANNEALING
MN AB-3 677 23 584 AFTER HOT-ROLLED SHEET ANYEALING
AC-3 1162 0.7 701 36 488 BEFORE HOT-ROLLED SHEET ANNEALING
AE-3 1050 1.3 623 49 525 AFTER HOT-ROLLED SHEET ANNEALING
AF-3 1029 1.4 605 76 604 AFTER HOT-ROLLED SHEET ANNEALING
AL-3 957 1.2 741 34 -715 BEFORE HOT-ROLLED SHEET ANNEALING
AM-3 1065 2.8 787 61 477 BEFORE HOT-ROLLED SHEET ANNEALING
AN-3 I 1018 1.4 721 87 1 646 I AFTER HOT-ROLLED SHEET ANNEALmG
AR-3 I 1207 I 1.8 I 626 1 82 1 480 I AFTER HOT-ROLLED SHEET ANNEALING
..\y.: ! : 5 - 6;: ,- 45-3 t3l',1bOl \ ' - 3 1 : - I ( . 8 833 1C i43 BCIOKI. IIOI-KOLLI:[) 5111.1:l :\\S!.:\!.l\(l
AW-3 1200 2.7 864 1 39 1 361 I AFTER HOT-ROLLED SHEET ANNEALING
AX-3 1191 2.1 791 1 37 1 553 I AFTER HOT-ROLLED SHEET ANNEALING
BC-3 I 1117 1 1.2 I 5-94 1 17 1 567 I AFTER HOT-ROLLED SHEET ANNEALING
BD-3 I 1145 1 1 .0 I 946 1 53 1 549 I AFTER HOT-ROLLED SHEET AhWEALING
BI-3 1027 2.5 946 1 36 1 -343 I AFTER HOT-ROLLED SHEET AhWEALING
BI-3 I I I 1038 1 83 1 389 I AFTER HOT-ROLLED SHEET ANNEALING
BN-3 I 1020 0.8 934 1 12 1 454 / AFTER HOT-ROLLED SHEET ANNEALING
BT-1 I I000 I 0.7 I 630 494 I BEFORE HOT-ROLLED SHEET ANNEALING
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES
. ---~, ~ ~ ~
("Chr)
No
--
COLD
ROLLING
REDUCTION
("C)
NOTE
HOT-ROLLED SHEET ANNEALING CONDITIONS
AA-3
AB-3
AC-3
AE-3
AF-3
AM-3
HEATING
RATE
UNTILTI
("Ckr) I ("C) I Olr)
135.8
AR-3
AS-3
AV-3
(%I
40.0 1 463.8 1 7.6
112.2
43.9
142.0
507.9
AL-3 / 31.8 1 491.6
AN-3 137.0 474.0 9.9 36.6 ~ 21.2 COMPARATIVE EXAMPLE
144.5 1 454.4 1 6.1 1 62.1 1 728 1 97.0
AW-3
AX-3
BC-3
lST COIL
SOAKING
TEMP.Tl'
5.5 / 72.4 1 684 1 38.6 68.6
BD-3
BF-3
BH-3
RI-3
* UNDERLINED BOLD NUMERICAL VfiUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES
54.4 1 732 1 65.3
537.0
461.4
530.3
INVENTION EXAMPLE
75.6
128.3
99.1
113.3
70.3
103.9
88.8
BK-3
BM-3
- BN-3
BT-3
RETENTIO
N TIME AT
T1
7.1
INVENTION EXAMPLE
8.2
6.2
9.2
505.9
490.5
492.1
59.3
78.3
128.8
49.4
17.0
6.2
5.2
8.2
51.8 726
506.5
525.1
491.5
74.1
102.6
54.8
137.2
HEATING
RATE
UNTIL T2
INVENTION EXAMPLE
30.5
36.2
52.3
474.5
483.0
545.1
498.5
63.0
67.6
39.3
7.9
8.2
5.6
7.5
zNDC OIL
SOAKING
TEMP. T2
751
675
711
5.8
6.0
7.0
9.6
497.8
461.3
500.0
490.0
RETENTION
TIME AT T2
710
711
734
41.6
46.8
16.9
16.1
68.7
7.9
71.0
16.8
6.7
9.1
2.8
4.7
COMPARATlVE EXAMPLE
66.8
71.9
17.7
30.4
19.4
73.3
36.0
31.0
6.0
5.1
716
681
747
749
763
72 1
716
26.2
69.4
17.9
674
757
727
72 1
54.5
-8.7
22.4
INVERTION EXAMPLE
COMPARATIVE EXAMPLE
COMPARATIVE EXAMPLE
96.3
59.8
78.1
8.7
59.1
14.1
99.6
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
INVENTION EXAMPLE
73.5
8.0
54.8
80.3
40.5
46.2
60.4
56.1
71.7
75.8
65.9
INVERTION EXAMPLE
INVEI'TION EXAMPLE
COMPARATIVE EXAMPLE
-86.1
66.1
-6.3
33.5
INVENTION EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
INVENTION EXAMPLE
COMPARATIVE EXAMPLE
COMPARATlVE EXAMPLE
[Tablc 8-1-11
COLD-ROLLED SHEET ANNEALING
I I I I
657 409.7
2-2 7.9 663
* UNDERLINED UOLII NUMERlCAL VALUES
INDICATE EXAMPLES OUTSIDE THE INVENTION
RANGES.
[Table 8- 1-21
[Tablc 8-1-21
* UNDERLWED BOLD NUMERICAL VALUES INIIICATE EXAMPLES OTITSIDE
THE INVENTION RANGES.
[0123]
[Table 8-1-31
[Table 8-1-31
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE TIE INVENTION IWGES.
[Table 8-2-11
[Table 8-2-11
I I COLD-ROLLED SHEET ANNEALING I
INDICATE EXAMPLES OUTSIDE THE INVENTION
RANGES.
[Table 8-2-21
[Table 8-2-21
* UNDERLINED BOLD NUMERlCAL VALUES lNDlCATE EXAMPLES OUTSIDE
THE INVENTION RANGES.
[0126]
[Table 8-2-31
[Table 8-2-31
NOTE
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE THE INVEN'I'ION RANGES.
[Table 9-1-11
INDICATE EXAMPLES OUTSIDE THE
INVENTION RANGES.
[Table 9-1-21
[Table 9-1 -21
* UNDERLNED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSlDE
THE INVENTION RANGES.
[0129]
[Table 9-1-31
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE THE WENTION RANGES.
[0130]
[Table 9-2-11
BT-3 1 22.9 1 702 1 1447 1 17.8 1 274.7
* UNDERLINED BOLD NUMERICAL VALUES
INDICATE EXAMPLES OUTSIDE 'THE
INVENTlON RANGES.
[Table 9-2-21
[Table 9-2-21
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE
THE INVENTION RANGES.
[Table 9-2-31
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES
OUTSIDE THE INVENTION RANGES.
FIG. 1 shows the shape of a test specimen used for evaluating the reduction in
area of a steel sheet during distortion at a high strain rate. The parallel portion in the
test specimen was 1.5 mm, the test specimen was pulled apart at a strolce rate of 900
mmlminute, the test specimen was broken, and the reduction in area of thc steel sheet
was obtained from a change in the sheet thickness at the center of the parallel portion
before and after a test.
[0134]
FIG. 2 shows the structure of Example U-1 in which ferrite and carbides were
made visible by etching a sample for which distortion at a high strain rate was stopped
at an elongation percentage of 13.4% using a 3% nitric acid-alcohol solution. It was
clear that cracking of carbides initiates from crystal interfaces present in carbide
particles.
[0135]
FIG. 3 shows a relationship between the reduction in area during distortion at
a high strain rate and the proportion of the number of carbides including a crystal
interface in each carbide particle to the number of all carbides regarding the invention
examples and the comparative examples in Tables 2-1 and 2-2 and the invention
examples and the comparative examples in Tables 5-1 to 5-6,6,7,8, and 9. It was
found that, when the composition is adjusted to be in the scope of the invention and the
proportion of the number of carbides including crystal interfaces was set to bc equal to
or lower than 20%, the reduction in area significantly improved.

I [Document Type] CLAIMS
1. A middlethigh carbon steel sheet, wherein
composition thereof comprises, by mass%:
C: 0.10% to 1.50%,
Si: 0.01% to 1.00%,
Mn: 0.01% to 3.00%,
P: 0.0001% to 0.1000%,
S: 0.0001% to 0.1000% and
a remainder consisting of Fe and impurities,
wherein the steel sheet has a structure in which a total volume percentage of a
martensite, a bainite, a pearlite, and a residual austenite is equal to or lower than 5.0%,
and a remainder thereof is a ferrite and carbides,
wherein a spheroidizing ratio of carbide particles is 70% to 9994, and
wherein a proportion of a number of the carbide particles including a crystal
interface at which an orientation difference is equal to or greater than 5" in the carbide
particles is equal to or lower than 20% of a total number of the carbide particles.
2. The middlehigh carbon steel sheet according to Claim 1,
wherein the composition of the steel sheet further includes, by mass%, one or
more selected from the group consisting of
Al: 0.001% to 0.500%,
N: 0.0001% to 0.0500%,
0: 0.0001% to 0.0500%,
Cr: 0.001% to 2.000%,
Mo: 0.001% to 2 .000%,
Ni: 0.001% to 2.000%,
Cu: 0.001% to 1.000%,
Nb: 0.001% to 1.000%,
V: 0.001% to 1.000%,
Ti: 0.001% to 1.000%,
B: 0.0001% to 0.0500%,
W: 0.001% to 1.000%,
Ta: 0.001% to 1.000%,
Sn: 0.001% to 0.020%,
Sb: 0.001% to 0.020%,
As: 0.001% to 0.020%,
Mg: 0.0001% to 0.0200%,
Ca: 0.001% to 0.020%,
Y: 0.001% to 0.020%,
Zr: 0.001% to 0.020%,
La: 0.001% to 0.020%, and
Ce: 0.001% to 0.020%.
3. A middlethigh carbon steel sheet and a method for manufacturing the same,
::
i;
wherein, when a billet having the composition according to Claim 1 or 2 is
directly hot-rolled or is temporary cooled, heated, and hot-rolled, finish hot-rolling is
completed in a temperature region of 600°C to l,OOO°C,
the hot-rolled steel sheet coiled at 350°C to 700°C is box-annealed,
cold-rolling of 10% io 80% is carried out, aiicl then cold-rolled-sheet.
annealing is carried out at an aiu~ealingte mperature of 650°C to 780°C for. a retelltion
time of 30 to 1,800 seconds in a continuous annealing line.