[Technical Field of the Invention]
[0001] The present invention relates to a middle/high 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] Middle/high 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 middle/high 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 /sec, 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 middle/high carbon steel sheets (for example, refer to Patent
Documents 1 to 6).
[0004] For example, Patent Document 1 discloses an invention of a method for manufacturing
a middle/high 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 µm to 1.50 µm in at least a final rolling path under conditions
of the 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.
[0005] Furthermore, Patent 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 µm, 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 µm 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 middle/high 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 µm, the spheroidizing ratio is equal to or higher than 90%,
and a texture satisfies an expression "(222)/(200)≥6-8.0×C (%)".
[0008] 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%, B: 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 (Ac1-10)°C, is cold-rolled at
a rolling reduction in a range of 40% to 80%, and is box-annealed at 650°C to (Ac1-10)°C.
[0010] 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]
[0011]
[Patent Document 1] Japanese Unexamined Patent Application, First Publication No.
2003-293042
[Patent Document 2] Japanese Unexamined Patent Application, First Publication No.
2003-147485
[Patent Document 3] Japanese Unexamined Patent Application, First Publication No.
2002-155339
[Patent Document 4] Japanese Unexamined Patent Application, First Publication No.
2000-328172
[Patent Document 5] Japanese Unexamined Patent Application, First Publication No.
H06-108158
[Patent Document 6] Japanese Unexamined Patent Application, First Publication No.
H11-61272
[Disclosure of the Invention]
[Problems to be Solved by the Invention]
[0012] The present invention has been made in consideration of the above-described circumstances,
and an object of the present invention is to provide a middle/high 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 middle/high 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 difficult to 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 like, and the steel sheets can only be manufactured
by achieving 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.
- (1) An aspect of the present invention provides a middle/high 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.
- (2) The middle/high carbon steel sheet according to (1), in which the 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%, O: 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%.
- (3) Another aspect of the present invention provides a method for manufacturing a
middle/high 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 1,000°C, the hot-rolled
steel sheet coiled at 350°C to 700° is box-annealed, cold-rolling 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]
[0016] According to the present invention, it is possible to provide a middle/high 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)]
[0017]
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]
[0018] Hereinafter, the present embodiment will be described in detail.
[0019] First, the reasons for limiting the chemical composition of a steel sheet according
to the present embodiment will be described. Here, regarding the composition, "%"
represents "mass%".
(C: 0.10% to 1.50%)
[0020] C is an element that increases the strength of steel by a heat treatment of quenching.
Middle/high 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 limit 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%.
(Si: 0.01% to 1.00%)
[0021] 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-rolled-sheet-annealing.
In a process of the Ostwald growth of carbide particles during cold-rolled-sheet-annealing,
when two or more particles that are adjacent to each other come into contact with
each other, crystal interfaces are introduced into the carbide particles. During the
distortion of the steel sheet, the crystal interfaces in the carbide particles serve
as the starting points of cracks. In order to suppress the above-described phenomenon,
it is necessary to decrease the growth rate of 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 cannot
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%.
(Mn: 0.01% to 3.00%)
[0022] 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 upper 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%.
(P: 0.0001% to 0.1000%)
[0023] 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%.
(S: 0.0001% to 0.1000%)
[0024] 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%.
[0025] 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 be essentially included, and thus the lower limit values of the elements described
below are 0%.
(Al: Preferably 0.001% to 0.500%)
[0026] Al is an element that serves as a deoxidizing agent of steel. When the Al content
is lower than 0.001%, effects of the inclusion of Al cannot be sufficiently obtained,
and thus the lower limit of the Al content may be set to 0.001%. On the other hand,
when the Al 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 Al content may be set to 0.500%. The Al content is more preferably
0.005% to 0.300% and still more preferably 0.010% to 0.100%.
(N: Preferably 0.0001% to 0.0500%)
[0027] 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%.
(O: Preferably 0.0001% to 0.0500%)
[0028] O is an element that accelerates the formation of coarse oxides in steel when a large
amount of O is included, and thus the O content is preferably lower. However, a decrease
in the O content to lower than 0.0001% leads to an increase in refining costs, and
thus the lower limit of the O content may be set to 0.0001%. On the other hand, when
the O 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 O content may be set to 0.0500%. The O content is
more preferably 0.0005% to 0.0250% and still more preferably 0.0010% to 0.0100%.
(Cr: Preferably 0.001% to 2.000%)
[0029] 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-sheet-annealing.
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 be 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-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 Cr content may be set to 2.000%. The Cr content is more preferably 0.005% to
1.500% and still more preferably 0.010% to 1.300%.
(Mo: Preferably 0.001% to 2.000%)
[0030] Mo is an element that, similar to Si, Mn, and Cr, suppresses the coarsening and joining
of carbide particles during hot-rolled-sheet-annealing and cold-rolled-sheet-annealing.
When the Mo content is lower than 0.001%, the above-described effect cannot be obtained,
and thus the lower limit of the 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-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 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%.
(Ni: Preferably 0.001% to 2.000%)
[0031] Ni is an element 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 be set to 2.000%. The Ni
content is more preferably 0.005% to 1.500% and still more preferably 0.005% to 0.700%.
(Cu: Preferably 0.001% to 1.000%)
[0032] 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 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 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%.
(Nb: Preferably 0.001% to 1.000%)
[0033] Nb is an element that forms carbonitrides and suppresses the coarsening and joining
of carbide particles during hot-rolled-sheet-annealing and cold-rolled-sheet-annealing.
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-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 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%.
(V: Preferably 0.001% to 1.000%)
[0034] V is also an element 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 above-described
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-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 V content 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%.
(Ti: Preferably 0.001% to 1.000%)
[0035] 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-sheet-annealing
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-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 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%.
(B: Preferably 0.0001% to 0.0500%)
[0036] 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 compounds are 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%.
(W: Preferably 0.001% to 1.000%)
[0037] 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-sheet-annealing
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-sheet-annealing 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%.
(Ta: Preferably 0.001% to 1.000%)
[0038] 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-rolled-sheet-annealing
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-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 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%.
(Sn: Preferably 0.001% to 0.020%)
[0039] 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 0.010%.
(Sb: Preferably 0.001 % to 0.020%)
[0040] 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 0.011%.
(As: Preferably 0.001% to 0.020%)
[0041] 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%.
(Mg: Preferably 0.0001% to 0.0200%)
[0042] 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%.
(Ca: Preferably 0.001% to 0.020%)
[0043] 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%.
(Y: Preferably 0.001% to 0.020%)
[0044] 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%.
(Zr: Preferably 0.001% to 0.020%)
[0045] 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 necessary. 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 0.015% and still more preferably equal to or lower than 0.010%.
(La: Preferably 0.001% to 0.020%)
[0046] 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%.
(Ce: Preferably 0.001% to 0.020%)
[0047] 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 reduction 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%.
[0048] Meanwhile, in the steel sheet according to the present embodiment, the remainder
of the composition described above is Fe and impurities.
[0049] 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 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
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.
[0050] Steel according to the present embodiment has a structure of substantially ferrite
and carbides. Meanwhile, carbides refer to not only cementite (Fe
3C) 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
(M
23C
6, M
6C, MC; here, M represents Fe and other alloy elements). Martensite, bainite, pearlite,
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 martensite, bainite, pearlite, and residual
austenite may be set to 0.0%.
[0051] 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.
[0052] 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%.
[0053] 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
mainly initiates from crystal interfaces at which a crystal orientation difference
is equal to or greater than 5° which are present in the carbides which have been considered
as a 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, 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 be 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 carbide 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%.
[0054] Subsequently, a method for observing and measuring the structure regulated above
will be described.
[0055] Ferrite, carbides, martensite, 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 µm, thereby finishing the observed sections to be mirror-like surfaces.
Next, the observed sections are etched using a 3% nitric acid-alcohol 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 µm×40
µm 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 µm, 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.
[0056] 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%.
[0057] 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 (311)
of austenite and of (200) and (211) of ferrite.
[0058] 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 µm and then carrying 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 the 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 diameter"=2×("area"/3.14)
1/2) of the respective carbides are obtained from the areas of the respective carbides,
and the 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 µm
2 are excluded from evaluation subjects.
[0059] A preferred range of the carbide particle diameter is 0.30 µm to 1.50 µm. In a case
in which the carbide particle diameter is smaller than 0.30 µm, the ferrite grain
size becomes too small, and thus the lower limit of the carbide particle diameter
is set to 0.30 µm. When the carbide particle diameter exceeds 1.50 µm, 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 µm. 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 (cementite and the like).
[0060] 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 then 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 SEM-EBSD 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 µm in a region 100 µm in the sheet thickness direction and 100 µm 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.
[0061] When the ferrite grain size in the structure after cold-rolled-sheet-annealing is
5 µm to 60 µm, 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 µm,
deformability is degraded, and thus the lower limit of the ferrite grain size is set
to 5 µm. In addition, when the ferrite grain size exceeds 60 µm, 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 limit of the ferrite grain size is set to equal to or
smaller than 60 µm. The ferrite grain size is measured by finishing the observation
surface to be a mirror-like surface by polishing in the above-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 µm to 50 µm.
[0062] Subsequently, a method for measuring reduction in area during distortion at a high
strain rate will be described.
[0063] In order to distort the steel sheet at a strain rate of 10 mm/sec 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 mm/minute, it becomes possible to impart strain to the parallel portions
in the test specimen at a strain rate extremely close to 10 mm/sec. 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, necking 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 necessary to set the ratio of the width of the parallel portion to the thickness
of the parallel portion to be 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 after tensile breakage using Equation (1)
[0064] 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 after 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 pieces due to breakage are respectively
measured, and the average of the measurement values at six points is used as the thickness
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".
[0065] Next, a method for manufacturing a steel sheet according to the present embodiment
will be described.
[0066] 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
annealing using a material having the above-described composition ranges.
[0067] The characteristics of a specific method for manufacturing a steel sheet according
to the present embodiment will be described below.
[0068] 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 1,000°C. The finishing-rolled steel strip is cooled on a run-out table (ROT)
at a cooling rate of 10 °C/second 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 hot-rolled 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 middle/high carbon steel sheet exhibiting excellent
reduction in area during distortion at a high strain rate.
[0069] Hereinafter, the method for manufacturing a steel sheet according to the present
embodiment will be specifically described.
(Hot-rolling)
[0070] When a slab (billet) having predetermined 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 1,000°C, and the obtained
steel strip is coiled in a temperature range of 350°C or more and less than 700°C.
[0071] 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 1,250°C, 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 950°C, and the lower limit of the heating
time is set to equal to or longer than 0.5 hours.
[0072] Finish hot-rolling is preferably ended at 600°C to 1,000°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 hot-rolling
temperature is set to equal to or higher than 600°C. In addition, when the finish
hot-rolling temperature exceeds 1,000°C, thick scales are generated on the steel sheet
during the passing of the 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 1,000°C, segregation of alloy elements such as Si and Mn into austenite grain
boundaries after the finish hot-rolling is accelerated, and the concentrations of
the alloy elements in austenite grains decrease, and thus carbides agglomerate during
hot-rolled-sheet-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 1,000°C.
[0073] The cooling rate of the steel strip on ROT after the finish hot-rolling is set to
10 °C/second to 100 °C/second. In a case in which the cooling rate is slower than
10 °C/second, the cooling rate is slow, and thus the growth of ferrite is accelerated,
and a structure in which ferrite, pearlite, 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 of the steel sheet, and thus the cooling rate is set to equal to
or faster than 10 °C/second. In addition, when the steel strip is cooled at a cooling
rate exceeding 100 °C/second 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 the above-described
low temperature transformation structures. 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 °C/second. 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 from the start of water injection to coiling in which a coiling device
is used.
[0074] The coiling temperature is set to 350°C to 700°C. When the coiling temperature is
lower than 350°C, 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 the
coiling temperature is set to be equal to or higher than 350°C. In addition, when
the coiling temperature exceeds 700°C, untransformed austenite transforms to pearlite
having coarse lamellar, and bulky needle-like cementite remains even after the cold-rolled-sheet-annealing,
and thus reduction in area is decreased. Therefore, the coiling temperature is set
to be equal to or lower than 700°C.
[0075] 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.
[0076] 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 °C/hr 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.
[0077] 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 even 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 retention time of box-annealing is set to equal to or
shorter than 100 hours. The lower limit of the retention time of box-annealing is
preferably two hours and more preferably five hours, and the upper limit thereof is
preferably 70 hours and more preferably 38 hours.
[0078] 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.
[0079] 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-rolled-sheet-annealing,
the coil after hot-rolled-sheet-annealing, which has been subjected to pickling before
or after the hot-rolled-sheet-annealing, is cold-rolled at a cold-rolling reduction
of 10% to 80%. In a case in which the cold-rolling reduction is lower than 10%, 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%. In addition, when the cold-rolling
reduction exceeds 80%, 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%.
[0080] 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 middle/high carbon steel sheet exhibiting excellent reduction in area during distortion
at a high strain rate.
[0081] 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-sheet-annealing,
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.
[0082] 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 650°C, 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 austenite being generated during the annealing
excessively increases, and thus it is not possible to suppress the generation of martensite,
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 of the 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-rolled-sheet-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 °C/second to 35 °C/second,
a cooling rate of 1 °C/second to 30 °C/second, 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.
[0083] According to the above-described method for manufacturing a steel sheet of the present
embodiment, it is possible to obtain a middle/high 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 residual 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
[0084] Next, the effects of the present invention will be described using examples.
[0085] 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.
[0086] 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 obtaining 250 mm-thick
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 36°C 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 °C/second, and were coiled at 510°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 °C/hour, 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 °C/hour, 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.
[0087] 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-1, 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-1, 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 the number of carbide particles including a crystal interface
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.
[0088] 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, but the C content was low, and high-strengthening was not possible
in the quenching step for producing products, and thus the steel sheet was evaluated
as fail. In Comparative Example K-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 O-1, the Mn content was 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 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 transformation 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 1]
[Table 1]
No |
C |
Si |
Mn |
P |
S |
NOTE |
A |
0.07 |
0.32 |
0.94 |
0.0162 |
0.0087 |
COMPARATIVE EXAMPLE |
B |
0.11 |
0.16 |
0.11 |
0.0048 |
0.0021 |
INVENTION EXAMPLE |
C |
0.20 |
0.41 |
2.22 |
0.0883 |
0.0027 |
INVENTION EXAMPLE |
D |
0.21 |
0.39 |
1.22 |
0.0064 |
0.0011 |
INVENTION EXAMPLE |
E |
0.29 |
0.26 |
0.82 |
0.0497 |
0.0039 |
INVENTION EXAMPLE |
F |
0.30 |
0.03 |
0.41 |
0.0089 |
0.0081 |
INVENTION EXAMPLE |
G |
0.32 |
0.22 |
0.22 |
0.0070 |
0.0213 |
INVENTION EXAMPLE |
H |
0.33 |
0.22 |
0.71 |
0.0088 |
0.0077 |
INVENTION EXAMPLE |
I |
0.35 |
0.31 |
1.33 |
0.0003 |
0.0049 |
INVENTION EXAMPLE |
J |
0.39 |
0.65 |
1.45 |
0.0012 |
0.0054 |
INVENTION EXAMPLE |
K |
0.41 |
0.38 |
0.004 |
0.0005 |
0.0061 |
COMPARATIVE EXAMPLE |
L |
0.43 |
0.34 |
0.67 |
0.1037 |
0.0082 |
COMPARATIVE EXAMPLE |
M |
0.45 |
0.26 |
0.03 |
0.0145 |
0.0138 |
INVENTION EXAMPLE |
N |
0.51 |
0.83 |
2.85 |
0.0163 |
0.0063 |
INVENTION EXAMPLE |
O |
0.51 |
0.68 |
3.16 |
0.0181 |
0.0081 |
COMPARATIVE EXAMPLE |
P |
0.54 |
0.18 |
0.36 |
0.0098 |
0.0082 |
INVENTION EXAMPLE |
Q |
0.59 |
0.20 |
0.67 |
0.0136 |
0.0054 |
INVENTION EXAMPLE |
R |
0.65 |
0.27 |
1.06 |
0.0107 |
0.0032 |
INVENTION EXAMPLE |
S |
0.69 |
0.24 |
0.77 |
0.0203 |
0.0756 |
INVENTION EXAMPLE |
T |
0.74 |
0.008 |
0.77 |
0.0112 |
0.0062 |
COMPARATIVE EXAMPLE |
U |
0.78 |
0.44 |
2.03 |
0.0266 |
0.0066 |
INVENTION EXAMPLE |
V |
0.82 |
0.47 |
0.87 |
0.0051 |
0.1134 |
COMPARATIVE EXAMPLE |
W |
0.98 |
1.09 |
2.11 |
0.0141 |
0.0146 |
COMPARATIVE EXAMPLE |
X |
0.92 |
0.35 |
1.86 |
0.0122 |
0.0016 |
INVENTION EXAMPLE |
Y |
0.94 |
0.54 |
1.92 |
0.0222 |
0.0449 |
INVENTION EXAMPLE |
Z |
1.02 |
0.09 |
0.88 |
0.0027 |
0.0003 |
INVENTION EXAMPLE |
AA |
1.12 |
0.76 |
1.39 |
0.0155 |
0.0098 |
INVENTION EXAMPLE |
AB |
1.33 |
0.50 |
0.78 |
0.0116 |
0.0273 |
INVENTION EXAMPLE |
AC |
1.44 |
0.14 |
0.31 |
0.0230 |
0.0168 |
INVENTION EXAMPLE |
AD |
1.58 |
0.27 |
0.51 |
0.0101 |
0.0099 |
COMPARATIVE EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 2-1]
[Table 2-2]
[0089] 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 thickness 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-hot-rolled at
870°C, then, were cooled to 510°C on ROT at a cooling rate of 45 °C/second, 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 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 °C/hour,
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 °C/hour, 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 out, and temper rolling
was carried out at a rolling reduction of 1.0%, thereby producing samples for characteristic
evaluation.
[0090] 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-1, AM-1, AN-1, AR-1, AS-1, AV-1, AW-1,
AX-1, BC-1, BD-1, BF-1, BH-1, BI-1, BJ-1, BK-1, BM-1, BN-1, 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.
[0091] In contrast, in Comparative Examples AG-1, AH-1, AO-1, AT-1, AU-1, 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-1,
AJ-1, AK-1, AQ-1, BE-1, BG-1, BL-1, BQ-1, and BR-1, 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 ferrite and carbides
increased, whereby stress accumulated in ferrite grain boundaries, and the reduction
in area was decreased. In Comparative Example AY-1, the O content was high, coarse
oxides were formed in steel, and fissures were initiated and propagated from the coarse
oxides as starting 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-B-carbides 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.
[Table 3-1]
[Table 3-1]
No |
C |
Si |
Mn |
P |
S |
Al |
N |
O |
Ti |
Cr |
Mo |
B |
Nb |
V |
NOTE |
AE |
0.16 |
0.29 |
0.53 |
0.0009 |
0.0174 |
|
|
|
|
|
|
0.0435 |
|
|
INVENTION EXAMPLE |
AF |
0.23 |
0.50 |
1.45 |
0.0528 |
0.0285 |
0.003 |
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
AG |
0.25 |
0.18 |
1.53 |
0.0046 |
0.0956 |
|
0.0052 |
|
|
1.69 |
|
|
|
|
COMPARATIVE EXAMPLE |
AH |
0.26 |
0.14 |
0.33 |
0.0199 |
0.0145 |
|
|
|
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
AI |
0.28 |
0.95 |
0.02 |
0.0715 |
0.0252 |
|
|
|
|
|
|
0.0376 |
1.053 |
|
COMPARATIVE EXAMPLE |
AJ |
0.30 |
0.41 |
2.13 |
0.0113 |
0.0177 |
|
|
|
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
AK |
0.30 |
0.15 |
1.05 |
0.0208 |
0.0811 |
|
0.0051 |
0.0042 |
1.072 |
|
|
|
0.902 |
|
COMPARATIVE EXAMPLE |
AL |
0.30 |
0.44 |
0.45 |
0.0766 |
0.0103 |
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
AM |
0.31 |
0.34 |
0.39 |
0.0067 |
0.0028 |
|
0.0048 |
|
|
|
|
|
0.598 |
|
INVENTION EXAMPLE |
AN |
0.31 |
0.18 |
1.14 |
0.0185 |
0.0023 |
|
|
0.0021 |
|
|
0.009 |
|
|
0.243 |
INVENTION EXAMPLE |
AO |
0.33 |
0.34 |
1.67 |
0.0053 |
0.0197 |
0.436 |
0.0010 |
|
|
1.02 |
|
0.0352 |
|
|
COMPARATIVE EXAMPLE |
AP |
0.38 |
0.83 |
2.43 |
0.0891 |
0.0224 |
|
0.0547 |
|
|
|
|
|
0.703 |
0.819 |
COMPARATIVE EXAMPLE |
AQ |
0.39 |
0.31 |
0.33 |
0.0009 |
0.0267 |
|
0.0076 |
|
0.958 |
0.36 |
|
|
|
0.690 |
COMPARATIVE EXAMPLE |
AR |
0.39 |
0.64 |
0.35 |
0.0041 |
0.0043 |
|
0.0113 |
|
0.003 |
|
|
|
|
|
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 3-2]
[Table 3-2]
No |
C |
Si |
Mn |
P |
S |
Al |
N |
O |
Ti |
Cr |
Mo |
B |
Nb |
V |
NOTE |
AS |
0.40 |
0.27 |
0.64 |
0.0062 |
0.0014 |
|
|
|
|
|
|
0.0271 |
|
0.750 |
INVENTION EXAMPLE |
AT |
0.43 |
0.24 |
2.62 |
0.0055 |
0.0097 |
0.556 |
|
|
|
|
|
0.0221 |
0.347 |
0.408 |
COMPARATIVE EXAMPLE |
AU |
0.44 |
0.36 |
0.40 |
0.0145 |
0.0018 |
|
|
0.0112 |
|
|
0.114 |
|
|
|
COMPARATIVE EXAMPLE |
AV |
0.44 |
0.07 |
1.17 |
0.0925 |
0.0215 |
|
|
|
0.138 |
|
|
|
|
|
INVENTION EXAMPLE |
AW |
0.45 |
0.11 |
0.73 |
0.0084 |
0.0009 |
0.356 |
|
|
|
|
0.872 |
0.0097 |
|
|
INVENTION EXAMPLE |
AX |
0.46 |
0.26 |
0.22 |
0.0165 |
0.0016 |
0.134 |
0.0074 |
|
|
0.012 |
|
0.0014 |
|
|
INVENTION EXAMPLE |
AY |
0.48 |
0.85 |
0.89 |
0.0061 |
0.0167 |
|
|
0.0520 |
|
|
1.962 |
0.024 |
|
|
COMPARATIVE EXAMPLE |
AZ |
0.48 |
0.89 |
2.10 |
0.0068 |
0.0185 |
0.446 |
|
|
0.215 |
|
|
|
0.043 |
0.053 |
COMPARATIVE EXAMPLE |
BA |
0.48 |
0.55 |
1.63 |
0.0048 |
0.0626 |
0.158 |
|
0.0017 |
0.103 |
|
|
0.0289 |
|
|
COMPARATIVE EXAMPLE |
BB |
0.51 |
0.42 |
2.94 |
0.0027 |
0.0217 |
|
0.0021 |
0.0354 |
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
BC |
0.51 |
0.98 |
1.32 |
0.0096 |
0.0085 |
|
|
|
|
|
|
|
0.009 |
|
INVENTION EXAMPLE |
BD |
0.52 |
0.69 |
0.94 |
0.0171 |
0.0489 |
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
BE |
0.55 |
0.58 |
1.31 |
0.0034 |
0.0085 |
|
0.0036 |
|
0.202 |
2.18 |
|
|
0.135 |
|
COMPARATIVE EXAMPLE |
BF |
0.58 |
0.24 |
0.28 |
0.0143 |
0.0089 |
|
|
0.0130 |
|
|
|
|
|
|
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 3-3]
[Table 3-3]
No |
C |
Si |
Mn |
P |
S |
A1 |
N |
O |
Ti |
Cr |
Mo |
B |
Nb |
V |
NOTE |
BG |
0.63 |
0.68 |
2.95 |
0.0025 |
0.0157 |
0.202 |
0.0014 |
|
|
|
2.121 |
0.0288 |
|
|
COMPARATIVE EXAMPLE |
BH |
0.65 |
0.18 |
0.71 |
0.0092 |
0.0038 |
0.024 |
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
BI |
0.66 |
0.03 |
0.52 |
0.0009 |
0.0068 |
|
|
|
0.491 |
|
|
|
|
|
INVENTION EXAMPLE |
BJ |
0.69 |
0.09 |
0.70 |
0.0181 |
0.0052 |
|
|
0.0320 |
|
1.02 |
|
|
|
|
INVENTION EXAMPLE |
BK |
0.70 |
0.69 |
2.58 |
0.0293 |
0.0914 |
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
BL |
0.71 |
0.09 |
0.49 |
0.0157 |
0.0071 |
|
|
|
|
|
|
|
0.102 |
1.034 |
COMPARATIVE EXAMPLE |
BM |
0.71 |
0.31 |
0.75 |
0.0122 |
0.0037 |
|
|
|
|
|
1.915 |
|
0.159 |
|
INVENTION EXAMPLE |
BN |
0.82 |
0.37 |
1.47 |
0.0067 |
0.0054 |
|
|
|
|
1.94 |
|
|
|
|
INVENTION EXAMPLE |
BO |
0.83 |
0.72 |
1.39 |
0.0432 |
0.0087 |
|
0.0185 |
|
|
0.25 |
|
|
0.032 |
|
COMPARATIVE EXAMPLE |
BP |
0.87 |
0.15 |
2.48 |
0.0298 |
0.0044 |
|
0.0198 |
|
|
|
|
0.0522 |
|
|
COMPARATIVE EXAMPLE |
BQ |
0.89 |
0.12 |
0.51 |
0.0188 |
0.0015 |
0.081 |
|
|
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
BR |
0.96 |
0.73 |
1.96 |
0.0239 |
0.0070 |
|
|
|
|
0.62 |
|
|
|
|
COMPARATIVE EXAMPLE |
BS |
1.21 |
0.34 |
0.75 |
0.0222 |
0.0222 |
|
|
|
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
BT |
1.32 |
0.91 |
1.81 |
0.0237 |
0.0686 |
|
0.0208 |
|
|
|
|
|
|
|
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 4-1]
[Table 4-1]
No |
Cu |
W |
Ta |
Ni |
Sn |
Sb |
As |
Mg |
Ca |
Y |
Zr |
La |
Ce |
NOTE |
AE |
|
|
|
|
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
AF |
|
|
|
|
|
|
|
0.0173 |
|
|
|
|
|
INVENTION EXAMPLE |
AG |
|
|
0.614 |
|
0.014 |
|
|
|
|
|
0.006 |
|
0.023 |
COMPARATIVE EXAMPLE |
AH |
|
|
|
|
|
|
|
|
0.024 |
|
|
|
|
COMPARATIVE EXAMPLE |
AI |
0.230 |
|
0.570 |
|
|
|
|
|
0.008 |
|
0.002 |
|
|
COMPARATIVE EXAMPLE |
AJ |
|
1.076 |
|
|
|
0.117 |
0.361 |
|
0.006 |
0.012 |
|
|
0.004 |
COMPARATIVE EXAMPLE |
AK |
|
|
|
|
0.007 |
|
|
|
|
|
|
0.013 |
|
COMPARATIVE EXAMPLE |
AL |
|
|
|
|
|
|
|
|
|
|
|
0.014 |
|
INVENTION EXAMPLE |
AM |
|
|
|
|
|
0.011 |
|
|
0.002 |
|
|
|
|
INVENTION EXAMPLE |
AN |
0.026 |
|
|
|
|
|
|
|
|
0.002 |
|
0.002 |
|
INVENTION EXAMPLE |
AO |
0.887 |
|
|
|
|
|
|
|
|
0.023 |
|
|
|
COMPARATIVE EXAMPLE |
AP |
|
|
|
|
|
|
|
|
|
0.020 |
|
0.011 |
0.017 |
COMPARATIVE EXAMPLE |
AQ |
|
|
|
2.107 |
|
|
|
|
|
|
0.014 |
|
|
COMPARATIVE EXAMPLE |
AR |
0.489 |
0.157 |
|
|
|
0.004 |
|
|
|
|
|
|
0.018 |
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 4-2]
[Table 4-2]
No |
Cu |
W |
Ta |
Ni |
Sn |
Sb |
As |
Mg |
Ca |
Y |
Zr |
La |
Ce |
NOTE |
AS |
|
|
|
0.659 |
0.013 |
|
|
|
0.017 |
0.013 |
|
|
|
INVENTION EXAMPLE |
AT |
0.847 |
0.014 |
|
|
|
|
0.206 |
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
AU |
0.098 |
|
0.694 |
|
|
|
|
0.0210 |
|
0.011 |
|
|
|
COMPARATIVE EXAMPLE |
AV |
|
|
|
|
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
AW |
|
|
0.144 |
1.464 |
|
|
0.006 |
|
|
|
|
|
|
INVENTION EXAMPLE |
AX |
|
|
|
|
|
|
|
|
|
0.019 |
|
|
0.003 |
INVENTION EXAMPLE |
AY |
|
0.505 |
|
1.495 |
|
|
|
|
|
|
|
0.005 |
|
COMPARATIVE EXAMPLE |
AZ |
|
|
|
|
|
|
0.024 |
|
|
|
|
0.012 |
|
COMPARATIVE EXAMPLE |
BA |
|
0.005 |
|
|
|
|
|
|
|
|
0.021 |
|
|
COMPARATIVE EXAMPLE |
BB |
|
0.753 |
0.297 |
|
0.022 |
|
|
|
|
|
0.017 |
|
|
COMPARATIVE EXAMPLE |
BC |
|
|
|
|
|
|
0.013 |
|
|
|
|
|
|
INVENTION EXAMPLE |
BD |
|
|
|
|
|
|
0.003 |
|
|
|
|
|
|
INVENTION EXAMPLE |
BE |
|
|
|
|
|
|
|
0.0091 |
|
|
|
0.018 |
|
COMPARATIVE EXAMPLE |
BF |
|
|
0.017 |
0.017 |
0.017 |
0.017 |
0.017 |
|
|
|
|
|
0.007 |
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 4-3]
[Table 4-3]
No |
Cu |
W |
Ta |
Ni |
Sn |
Sb |
As |
Mg |
Ca |
Y |
Zr |
La |
Ce |
NOTE |
BG |
0.294 |
|
|
|
|
|
0.018 |
|
|
|
|
|
|
COMPARATIVE EXAMPLE |
BH |
|
|
|
|
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
BI |
0.116 |
|
|
|
|
0.013 |
|
|
0.008 |
|
0.017 |
0.007 |
|
INVENTION EXAMPLE |
BJ |
|
|
|
|
0.006 |
|
|
|
|
|
|
0.018 |
|
INVENTION EXAMPLE |
BK |
|
0.434 |
|
|
|
|
|
|
|
|
|
|
|
INVENTION EXAMPLE |
BL |
|
|
0.165 |
|
|
0.002 |
0.004 |
|
|
0.015 |
|
|
|
COMPARATIVE EXAMPLE |
BM |
|
|
0.724 |
0.061 |
|
|
|
0.0074 |
|
|
0.002 |
|
|
INVENTION EXAMPLE |
BN |
|
|
|
|
|
|
|
|
0.010 |
|
0.013 |
|
|
INVENTION EXAMPLE |
BO |
|
|
|
|
|
0.021 |
|
|
|
0.006 |
0.002 |
|
|
COMPARATIVE EXAMPLE |
BP |
|
|
0.402 |
0.006 |
0.011 |
|
|
0.0042 |
|
|
|
|
|
COMPARATIVE EXAMPLE |
BQ |
1.072 |
|
|
|
0.002 |
|
0.014 |
|
|
|
|
0.014 |
0.002 |
COMPARATIVE EXAMPLE |
BR |
|
0.892 |
1.054 |
0.452 |
|
|
|
0.0124 |
|
0.005 |
|
|
|
COMPARATIVE EXAMPLE |
BS |
|
|
|
|
|
|
|
|
|
|
|
0.022 |
|
COMPARATIVE EXAMPLE |
BT |
|
|
|
|
|
|
|
|
|
0.008 |
|
|
|
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 5-1]
[Table 5-2]
[Table 5-3]
[Table 5-4]
[Table 5-5]
[Table 5-6]
[0092] 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 1, 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-rolled steel annealing, pickling, cold-rolling, and cold-rolled-sheet-annealing
were carried out, thereby manufacturing samples for characteristic evaluation.
[0093] 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, Z-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, bainite,
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.
[0094] 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, Q-3, and BI-3, the coiling temperatures were low, the structures of the hot-rolled
sheet were fine, and the structures after the cold-rolled-sheet-annealing were also
fine, and thus deformability degraded, and the reductions in area were decreased during
distortion at a high strain rate.
[0095] 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 cold-rolled-sheet-annealing
became fine, deformability degraded, and the reductions in area were decreased. In
Comparative Examples S-2, AW-2, AC-3, and BH-3, the cold-rolling reductions were low,
and thus the ferrite grain sizes after the cold-rolled-sheet-annealing 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-rolled-sheet-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 pearlitic 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 I-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 process, 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-rolled-sheet-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.
[Table 6-1-1]
[Table 6-1-2]
[Table 6-2-1]
[Table 6-2-2]
[Table 7-1-1]
[Table 7-1-1]
No |
HOT ROLLING CONDITIONS |
TIMING OF PICKLING |
HEATING TEMPERATURE (°C) |
HEATING TIME (hr) |
FINISHING TEMPERATURE (°C) |
ROT COOLING RATE (°C/s) |
COILING TEMPERATURE (°C) |
B-3 |
1112 |
0.9 |
793 |
14 |
445 |
BEFORE HOT-ROLLED SHEET ANNEALING |
C-3 |
- |
- |
1026 |
79 |
465 |
AFTER HOT-ROLLED SHEET ANNEALING |
D-3 |
1126 |
2.0 |
900 |
53 |
370 |
AFTER HOT-ROLLED SHEET ANNEALING |
E-3 |
1120 |
0.9 |
607 |
13 |
361 |
AFTER HOT-ROLLED SHEET ANNEALING |
F-3 |
1164 |
1.0 |
782 |
98 |
437 |
AFTER HOT-ROLLED SHEET ANNEALING |
G-3 |
1156 |
3.0 |
840 |
70 |
549 |
AFTER HOT-ROLLED SHEET ANNEALING |
H-3 |
987 |
1.5 |
759 |
52 |
645 |
BEFORE HOT-ROLLED SHEET ANNEALING |
1-3 |
963 |
0.6 |
682 |
25 |
356 |
AFTER HOT-ROLLED SHEET ANNEALING |
J-3 |
1214 |
1.8 |
629 |
42 |
551 |
AFTER HOT-ROLLED SHEET ANNEALING |
M-3 |
1059 |
1.1 |
847 |
69 |
598 |
AFTER HOT-ROLLED SHEET ANNEALING |
N-3 |
- |
- |
671 |
32 |
659 |
AFTER HOT-ROLLED SHEET ANNEALING |
P-3 |
1148 |
1.0 |
630 |
99 |
429 |
BEFORE HOT-ROLLED SHEET ANNEALING |
Q-3 |
1164 |
0.8 |
669 |
16 |
337 |
AFTER HOT-ROLLED SHEET ANNEALING |
R-3 |
1110 |
1.5 |
920 |
79 |
410 |
BEFORE HOT-ROLLED SHEET ANNEALING |
S-3 |
1030 |
2.0 |
876 |
16 |
381 |
BEFORE HOT-ROLLED SHEET ANNEALING |
U-3 |
1165 |
1.6 |
963 |
20 |
352 |
BEFORE HOT-ROLLED SHEET ANNEALING |
X-3 |
1060 |
2.9 |
577 |
56 |
423 |
AFTER HOT-ROLLED SHEET ANNEALING |
Y-3 |
1039 |
1.4 |
798 |
79 |
718 |
AFTER HOT-ROLLED SHEET ANNEALING |
Z-3 |
1239 |
2.1 |
809 |
57 |
650 |
AFTER HOT-ROLLED SHEET ANNEALING |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 7-1-2]
[Table 7-1-2]
No |
HOT-ROLLED SHEET ANNEALING CONDITIONS |
COLD ROLLING REDUCTION (%) |
NOTE |
HEATING RATE UNTIL T1 (°C/hr) |
1ST COIL SOAKING TEMP. T1 (°C) |
RETENTION TIME AT T1 (hr) |
HEATING RATE UNTIL T2 (°C/hr) |
2ND COIL SOAKING TEMP. T2 (°C) |
RETENTION TIME AT T2 (hr) |
B-3 |
90.6 |
537.2 |
9.2 |
78.2 |
711 |
97.2 |
44.0 |
COMPARATIVE EXAMPLE |
C-3 |
109.1 |
459.7 |
8.9 |
48.5 |
693 |
99.0 |
34.1 |
COMPARATIVE EXAMPLE |
D-3 |
149.3 |
533.6 |
5.9 |
32.3 |
690 |
18.3 |
56.5 |
COMPARATIVE EXAMPLE |
E-3 |
31.9 |
528.7 |
5.1 |
59.6 |
740 |
47.8 |
39.1 |
COMPARATIVE EXAMPLE |
F-3 |
45.0 |
523.7 |
5.2 |
49.4 |
715 |
29.6 |
79.9 |
INVENTION EXAMPLE |
G-3 |
100.6 |
459.3 |
2.7 |
58.7 |
755 |
70.1 |
13.8 |
INVENTION EXAMPLE |
H-3 |
74.3 |
476.2 |
8.3 |
12.6 |
715 |
59.0 |
64.8 |
INVENTION EXAMPLE |
I-3 |
117.4 |
473.0 |
7.6 |
12.5 |
702 |
47.5 |
54.6 |
INVENTION EXAMPLE |
J-3 |
102.8 |
477.2 |
8.7 |
63.6 |
728 |
41.2 |
88.7 |
COMPARATIVE EXAMPLE |
M-3 |
142.9 |
526.6 |
4.8 |
6.8 |
765 |
31.1 |
55.2 |
INVENTION EXAMPLE |
N-3 |
53.1 |
486.2 |
4.0 |
8.6 |
749 |
71.2 |
45.6 |
INVENTION EXAMPLE |
P-3 |
91.1 |
501.1 |
7.0 |
14.9 |
706 |
41.9 |
56.5 |
INVENTION EXAMPLE |
Q-3 |
39.7 |
499.7 |
3.9 |
32.0 |
715 |
92.7 |
40.0 |
COMPARATIVE EXAMPLE |
R-3 |
70.1 |
485.6 |
4.5 |
46.5 |
751 |
61.4 |
65.5 |
INVENTION EXAMPLE |
S-3 |
124.2 |
505.8 |
3.7 |
41.8 |
747 |
53.8 |
25.4 |
INVENTION EXAMPLE |
U-3 |
114.4 |
492.8 |
4.1 |
50.0 |
699 |
28.5 |
51.2 |
INVENTION EXAMPLE |
X-3 |
85.2 |
516.3 |
7.7 |
45.1 |
754 |
81.7 |
61.4 |
COMPARATIVE EXAMPLE |
Y-3 |
35.2 |
533.7 |
6.1 |
50.4 |
710 |
80.3 |
57.2 |
COMPARATIVE EXAMPLE |
Z-3 |
127.1 |
451.1 |
3.7 |
49.3 |
729 |
68.8 |
23.5 |
COMPARATIVE EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 7-2-1]
[Table 7-2-1]
No |
HOT ROLLING CONDITIONS |
TIMING OF PICKLING |
HEATING TEMPERATURE (°C) |
HEATING TIME (hr) |
FINISHING TEMPERATURE (°C) |
ROT COOLING RATE (°C/s) |
COILING TEMPERATURE (°C) |
AA-3 |
1183 |
0.5 |
805 |
85 |
654 |
AFTER HOT-ROLLED SHEET ANNEALING |
AB-3 |
- |
- |
677 |
23 |
584 |
AFTER HOT-ROLLED SHEET ANNEALING |
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 |
1018 |
1.4 |
721 |
87 |
646 |
AFTER HOT-ROLLED SHEET ANNEALING |
AR-3 |
1207 |
1.8 |
626 |
82 |
480 |
AFTER HOT-ROLLED SHEET ANNEALING |
AS-3 |
1150 |
2.0 |
634 |
54 |
458 |
BEFORE HOT-ROLLED SHEET ANNEALING |
AV-3 |
1171 |
0.8 |
833 |
40 |
542 |
BEFORE HOT-ROLLED SHEET ANNEALING |
AW-3 |
1200 |
2.7 |
864 |
39 |
361 |
AFTER HOT-ROLLED SHEET ANNEALING |
AX-3 |
1191 |
2.1 |
791 |
37 |
553 |
AFTER HOT-ROLLED SHEET ANNEALING |
BC-3 |
1117 |
1.2 |
594 |
17 |
567 |
AFTER HOT-ROLLED SHEET ANNEALING |
BD-3 |
1145 |
1.0 |
946 |
53 |
549 |
AFTER HOT-ROLLED SHEET ANNEALING |
BF-3 |
1158 |
1.9 |
850 |
24 |
580 |
AFTER HOT-ROLLED SHEET ANNEALING |
BH-3 |
1241 |
2.5 |
970 |
84 |
531 |
BEFORE HOT-ROLLED SHEET ANNEALING |
BI-3 |
1027 |
2.5 |
946 |
36 |
343 |
AFTER HOT-ROLLED SHEET ANNEALING |
BJ-3 |
- |
- |
1038 |
83 |
389 |
AFTER HOT-ROLLED SHEET ANNEALING |
BK-3 |
1038 |
1.6 |
877 |
92 |
604 |
AFTER HOT-ROLLED SHEET ANNEALING |
BM-3 |
1150 |
1.0 |
657 |
98 |
634 |
AFTER HOT-ROLLED SHEET ANNEALING |
BN-3 |
1020 |
0.8 |
934 |
12 |
454 |
AFTER HOT-ROLLED SHEET ANNEALING |
BT-3 |
1000 |
0.7 |
630 |
42 |
494 |
BEFORE HOT-ROLLED SHEET ANNEALING |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 7-2-2]
[Table 7-2-2]
No |
HOT-ROLLED SHEET ANNEALING CONDITIONS |
COLD ROLLING REDUCTION (%) |
NOTE |
HEATING RATE UNTIL T1 (°C/hr) |
1ST COIL SOAKING TEMP. T1 (°C) |
RETENTION TIME AT T1 (hr) |
HEATING RATE UNTIL T2 (°C/hr) |
2ND COIL SOAKING TEMP. T2 (°C) |
RETENTION TIME AT T2 (hr) |
AA-3 |
40.0 |
463.8 |
7.6 |
54.4 |
732 |
65.3 |
17.0 |
INVENTION EXAMPLE |
AB-3 |
112.2 |
537.0 |
6.2 |
63.0 |
710 |
31.0 |
54.5 |
INVENTION EXAMPLE |
AC-3 |
43.9 |
461.4 |
5.2 |
67.6 |
711 |
6.0 |
8.7 |
COMPARATIVE EXAMPLE |
AE-3 |
142.0 |
530.3 |
8.2 |
39.3 |
734 |
5.1 |
22.4 |
INVENTION EXAMPLE |
AF-3 |
135.8 |
507.9 |
5.5 |
72.4 |
684 |
38.6 |
68.6 |
INVENTION EXAMPLE |
AL-3 |
31.8 |
491.6 |
7.1 |
51.8 |
726 |
7.9 |
41.6 |
COMPARATIVE EXAMPLE |
AM-3 |
144.5 |
454.4 |
6.1 |
62.1 |
728 |
97.0 |
75.6 |
INVENTION EXAMPLE |
AN-3 |
137.0 |
474.0 |
9.9 |
36.6 |
770 |
63.6 |
21.2 |
COMPARATIVE EXAMPLE |
AR-3 |
128.3 |
505.9 |
8.2 |
30.5 |
751 |
66.8 |
26.2 |
INVENTION EXAMPLE |
AS-3 |
99.1 |
490.5 |
6.2 |
36.2 |
675 |
71.9 |
69.4 |
COMPARATIVE EXAMPLE |
AV-3 |
113.3 |
492.1 |
9.2 |
52.3 |
711 |
17.7 |
17.9 |
COMPARATIVE EXAMPLE |
AW-3 |
70.3 |
506.5 |
8.2 |
46.8 |
716 |
96.3 |
40.5 |
INVENTION EXAMPLE |
AX-3 |
103.9 |
525.1 |
5.6 |
16.9 |
681 |
59.8 |
46.2 |
INVENTION EXAMPLE |
BC-3 |
88.8 |
491.5 |
7.5 |
16.1 |
747 |
78.1 |
60.4 |
COMPARATIVE EXAMPLE |
BD-3 |
59.3 |
474.5 |
5.8 |
30.4 |
674 |
73.5 |
86.1 |
COMPARATIVE EXAMPLE |
BF-3 |
78.3 |
483.0 |
6.0 |
19.4 |
757 |
8.0 |
66.1 |
INVENTION EXAMPLE |
BH-3 |
128.8 |
545.1 |
7.0 |
73.3 |
727 |
54.8 |
6.3 |
COMPARATIVE EXAMPLE |
BI-3 |
49.4 |
498.5 |
9.6 |
36.0 |
721 |
80.3 |
33.5 |
COMPARATIVE EXAMPLE |
BJ-3 |
52.7 |
524.3 |
4.0 |
18.8 |
718 |
38.5 |
61.3 |
COMPARATIVE EXAMPLE |
BK-3 |
74.1 |
497.8 |
6.7 |
68.7 |
749 |
8.7 |
56.1 |
INVENTION EXAMPLE |
BM-3 |
102.6 |
461.3 |
9.1 |
7.9 |
763 |
59.1 |
71.7 |
INVENTION EXAMPLE |
BN-3 |
54.8 |
500.0 |
2.8 |
71.0 |
721 |
14.1 |
75.8 |
COMPARATIVE EXAMPLE |
BT-3 |
137.2 |
490.0 |
4.7 |
16.8 |
716 |
99.6 |
65.9 |
INVENTION EXAMPLE |
* UNDERLINED BOLD NUMERICAL VALUES INDICATE EXAMPLES OUTSIDE THE INVENTION RANGES. |
[Table 8-1-1]
[Table 8-1-2]
[Table 8-1-3]
[Table 8-2-1]
[Table 8-2-2]
[Table 8-2-3]
[Table 9-1-1]
[Table 9-1-2]
[Table 9-1-3]
[Table 9-2-1]
[Table 9-2-2]
[Table 9-2-3]
[0096] 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 stroke rate
of 900 mm/minute, the test specimen was broken, and the reduction in area of the steel
sheet was obtained from a change in the sheet thickness at the center of the parallel
portion before and after a test.
[0097] 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.
[0098] 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 be equal to or lower
than 20%, the reduction in area significantly improved.