TECHNICAL FIELD
The present invention relates to a thin cold-rolled
steel plate used in home appliances and the like, and a
preparation method thereof, and more particularly, to a thin
10 cold-rolled steel plate having high strength and high
formability, and a preparation method thereof.
Since most existing steel sheets used in home
appliances are low carbon steel sheets, the formability
thereof is considered to be an important factor while the
15 strength thereof may not be considered.
Particularly, since excellent deep drawing qualities
(EDDQ) or higher strength steel sheets requiring high degrees
of formability are rather concentrated on the formability
thereof, the strength of these steel sheets is controlled to
20 a level below a specific value.
However, with recent demands, such as for low
production costs, high fuel efficiency, slimness, and the
like in product families including automobiles, home
appliances, and the like, which mainly use cold-rolled steel
25 plates, the most important keywords are thin and high
strength. That is, since the use of thin steel plates may
decrease the overall weight of products in which the
ultrathin steel sheet is used, low production costs, thinner
products and a range of product designs may be realized.
30 Thus, thinning and high strength may provide the effect of
3
killing three birds with one stone.
A great deal of research into the development of an
thin steel sheet product having high strength and high
formability has been undertaken.
Such research may be generally classified into 5 o 1)
texture (transformation) strengthening using a transformation
generated during the manufacturing of steel plates, 2) solid
solution strengthening controlling a component that may be
solid-solutioned in the matrix phase of steel, 3)
10 precipitation strengthening used to improve the strength of a
pure metal through the distribution of precipitates, and 4)
work strengthening causing work hardening by performing
secondary rolling of a steel strip completely recrystallized
in a final annealing process.
15 Such related art inventions may be generally divided
into two types, 1) a double reducing (DR) type process which
uses secondary rolling according to the applied process, and
2) a DR-omitted process which does not use secondary rolling.
That is, the deformation strengthening, the solid solution
20 strengthening, the precipitation strengthening, and the like
may be also classified into a DR type process and a DRomitted
process according to the existence or absence of
secondary rolling.
Since the DR type process to improve strength using
25 secondary rolling generates defects such as dislocations in
steel, which are inevitably accompanied by an increase in
strength due to secondary rolling to slowly increase the
strength of a steel plate while causing a sharp drop in
elongation, secondary rolling may not be used on a portion of
30 a steel plate in which a high degree of formability will be
4
required.
For example, since most steel plates subject to
secondary rolling have a low elongation of 2-3%, such steel
sheets have the drawback of cracks generated in the rolling
direction through a reduction in formability due to the 5 low
elongation and influence of rolled grains generated in the
secondary rolling.
Carbon steels in these related art inventions may be
classified, based on carbon content, as ultra low carbon
10 steel having a carbon content not more than 0.01 wt%, low
carbon steel having carbon content of 0.01< wt% C <0.1,
medium carbon steel having carbon content of 0.1< wt% C <0.25,
and high carbon steel having carbon content of 0.25wt% or
more.
15 Ultra low carbon steel is mainly used for cans, and as
the ultra low carbon steel related art inventions, Japanese
Patent Application No. 1995-274558 in which a reduction ratio
in secondary reduction is reduced and Mn content is
controlled to improve strength, and Japanese Patent
20 Application No. 1997-216980 in which a reduction ratio is
controlled to improve workability are disclosed.
Also, Japanese Patent Application Nos. 2002-307898 and
2002-201574 disclose inventions in which high temperature
strengths of the same steel sheets are improved by using
25 solid solution strengthening and precipitation strengthening
of Mn, P, TiC, or the like. However, since ultra low carbon
steel has limited strength and the elongation of ultra low
carbon steel may be reduced to a very low level in the case
that secondary rolling is performed thereon in order to
30 improve the strength thereof, such ultra low carbon steel has
5
limitations in terms of high formability and in uses thereof
in high strength products.
Also, most high strength steel plates formed of low
carbon steels are used as black plates (BP) for cans, and
related art inventions regarding BPs for cans 5 include a
technique (Japanese Patent Application No. 1990-052642) in
which high nitrogen steel and a double reducing mill (DRM)
are used, a technique (Japanese Patent Application No. 1996-
239734) in which an Mn content is increased in a steel sheet
10 and continuous lubrication rolling and secondary rolling are
applied thereto, a technique (Japanese Patent Application No.
1997-040883) in which an overaging treatment effect is used,
and a technique (Japanese Patent Application No. 2006-074140)
in which rapid cooling is performed and a texture obtained
15 by rapid cooling is used.
However, these related art inventions have limitations
in that low carbon steels have low strength levels and high
cooling rates that are difficult to realize in general
continuous annealing processes, such annealing processes
20 being required even in the case that strength levels are high,
and the range of obtained elongation is lower than a target
elongation range.
Also, since most high carbon steels having a carbon
content of 0.2 wt% or more have difficulties in the reduction
25 of a sectional area due to initial high strengths in PCM and
difficulties in leveling for control of form after reduction,
they are not used for thin cold-rolled products.
Recently, steel plates have been developed in which
through combinations of these concepts, phosphorous (P) is
30 added to perform solid strengthening of matrix textures in
6
medium carbon steels, while at the same time the matrix
textures are formed to have a two phase texture of ferrite +
pearlite, and secondary rolling is controlled to a low
percentage of 10% or less to maximize the combination of
strength and elongation (Korean Patent Application No. 5 2009-
0084530).
Particularly, this Korean Patent discloses a method in
which all of solid solution strengthening, texture control,
and work hardening through secondary rolling are used to
10 obtain a higher strength level compared with other techniques
(Y.S.>650 MPa) and the amount of secondary rolling is low, so
as to obtain thin cold-rolled steel plates having superior
formability in the rolling direction.
However, these patents have limitations in that the
15 processes are complicated due to the use of secondary rolling,
and even when the amount of rolling is small, dislocations
are generated due to the effects of rolling to cause
differences in formability between the rolling direction and
the vertical direction.
20
DISCLOSURE OF THE INVENTION
TECHNICAL PROBLEM
An aspect of the present disclosure may provide a thin
cold-rolled steel plate having high strength and high
25 formability and a method of producing the same.
An aspect of the present disclosure may also provide a
method of producing a thin cold-rolled steel plate having
high strength and high formability without secondary rolling
by appropriately controlling compositions and manufacturing
30 conditions of steel.
7
TECHNICAL SOLUTION
According to an aspect of the present disclosure, a
thin cold-rolled steel plate, having high strength and high
formability, may include: 0.15-0.25 wt% of carbon (5 C), 1.5-
2.5 wt% of manganese (Mn), 0.1-1.0 wt% of silicon (Si), 0.01-
0.05 wt% of titanium (Ti), 5-30 ppm of boron (B), and a
balance of Fe and inevitable impurities, wherein the
microstructure comprises 70-100 vol% of bainite, and 0-30
10 vol% of ferrite.
According to another aspect of the present disclosure,
a method of producing a thin cold-rolled steel plate includes
heating a steel slab including 0.15-0.25 wt% of carbon (C),
1.5-2.5 wt% of manganese (Mn), 0.1-1.0 wt% of silicon (Si),
15 0.01-0.05 wt% of titanium (Ti), 5-30 ppm of boron (B), and a
balance of Fe and inevitable impurities; performing hotfinish
rolling of the steel slab at a temperature equal to an
Ar3 transformation point or above to form a hot-rolled steel
plate; coiling the hot-rolled steel plate at a temperature of
20 500-800°C, cold-rolling the hot-rolled steel plate at a
reduction ratio of 50-90% to form a cold-rolled steel plate;
maintaining the cold-rolled steel sheet in a continuous
annealing line at an annealing temperature of 750-850°C for
30 seconds or more; cooling the cold-rolled steel sheet at a
25 temperature of 250-450°C at a cooling rate of 10-50°C/sec;
maintaining the cooled steel plate at the temperature of 250-
450°C for 50 seconds or more; and recooling the steel plate.
ADVANTAGEOUS EFFECTS
30 As set forth above, according to exemplary embodiments
8
of the present disclosure, a thin steel sheet having high
strength and high formability may be provided, and thus can
be effectively used in high strength and thin cold-rolled
products and the like, requiring strength of 300 HV or higher
on the basis of HV 500g in addition to product 5 components
supporting the strength of a chassis such as components using
in products such as laptop computers, LCD monitors, LCD
devices, PMP or LED TVs, and the like.
10 BRIEF DESCRIPTION OF THE DRAWINGS
FIG. 1 provides optical photomicrographs of an
inventive steel and a comparative steel, and particularly,
FIG. 1A is an optical photomicrograph of an inventive steel
and FIG. 1B is an optical photomicrograph of a comparative
15 steel.
FIG. 2 provides scanning electron microscopic (SEM)
photographs of inventive steels taken at magnifications of
1000x (FIG. 2A), 2000x (FIG. 2B), and 5000x (FIG. 2C).
20 MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described in
detail.
In order to obtain a low temperature transformation
texture at a low cooling rate, relatively expensive alloying
25 elements such as Nb, Mo, Ti, and the like are excluded and
relatively inexpensive Mn, B, and the like are added in steel
to secure high hardenability through content control of Mn, B,
and the like, thereby allowing for a low temperature
transformation texture even at a slow cooling rate, for
30 example, 30°C/sec. or lower, the cooling rate in an annealing
9
process in a continuous annealing line (CAL).
Inventive steels are characterized by textures
including 70-100 vol.% of bainite and 0-30 vol.% of ferrite,
and since bainite may be obtained at a general cooling rate,
inventive steels have advantages of reduced distortion 5 and
superior workability as compared with martensite steels.
Also, since inventive steel sheets, not subject to
secondary rolling, have a hardness of HV 500, higher by HV
300 or more than high strength thin steel plates subject to
10 secondary rolling and have a hardness of HV 200-250, the
inventive steel sheets do not have anisotropic properties in
several rolling directions, generated during secondary
rolling.
Hereinafter, compositions of steels according to the
15 present disclosure will be described (unit: wt%).
Carbon (C) is preferably contained in an amount of
0.15% or more in order to control a texture of steel and to
secure sufficient strength in production of thin cold-rolled
steel plates, but the upper limit of carbon content is
20 limited to 0.25% in consideration of control of a
precipitated amount of carbides, workability of steel sheets,
possibility of cold-rolling, shape deterioration, and
reductions of threading in annealing.
Manganese (Mn) reduces an Ar3 transformation point
25 temperature, improves hardenability in cooling to delay the
formation of a transformed phase such as pearlite at a low
cooling rate such that bainite phases may be formed at a
typical cooling rate, and is an essential component, added in
order to prevent hot shortness due to sulfur (S) impurities.
30 Mn is preferably added in a content of 1.5% or more, but the
10
content of Mn is controlled to 2.5% or less in consideration
of cold-rollability, brittleness of slabs, and the like.
The Ar3 transformation point temperature is a reverse
transformation temperature to form an austenite pool which
allows a reverse transformation to occur upon cooling, 5 after
a continuous annealing process.
Boron (B) is a main element which improves
hardenability together with Mn to allow a bainite phase to be
formed even at a typical cooling rate during the annealing
10 heat treatment. When the content of B is below 5 ppm, the
effect of B cannot be expected, and when the content of B is
more than 30 ppm, boron-based precipitates may be formed
excessively at grain boundaries to negatively influence
physical properties of steel. Therefore, the content of B is
15 limited to 5-30 ppm.
Titanium (Ti) is an element added to allow the effect
of B to be more stably obtained, and is added as a scavenger
element for suppressing the formation of boron nitrides due
to the bonding of nitrogen (N) and boron remaining in steel.
20 Therefore, the content of Ti is determined in proportional to
the content of N remaining in steel and thus is limited to
0.01-0.05%.
Silicon (Si) is an element added as a deoxidizer and
for solid solution strengthening; however, when the content
25 of Si exceeds 1.0%, crack embrittlement occurs.
It is preferable that the product of contents of C, Mn
and B satisfy the relationship of 1.13*10-4 < wt% C * wt% Mn
* wt% B < 1.875*10-3.
When the product of contents is more than 1.875*10-3,
30 embrittlement may occur and rollability is reduced, and when
11
the product of contents is below 1.13*10-4, an Ar3
transformation point temperature rises and hardenability is
reduced so that bainite is not sufficiently formed.
Besides the above-described components, aluminum (Al),
phosphorous 5 (P) and sulfur (S) may be added.
Preferably, Al may be included in an amount up to 0.06%,
and P and S may be included in an amount up to 0.03%,
respectively.
Cold-rolled steel plates according to embodiments of
10 the present disclosure include 70-100 vol.% of bainite and 0-
30 vol.% of ferrite.
Since the bainite texture may be obtained even at a
typical cooling rate, the cold-rolled steel plates according
to embodiments of the present disclosure may have less
15 distortion as compared with martensite steel sheets, to thus
improve the workability and formability thereof.
The cold-rolled steel plates according to embodiments
of the present disclosure may include ferrite in an amount up
to 30 vol.%.
20 Ferrite is a microstructure allowing for the ductility
of steel sheets to be increased, and is included in an amount
up to 30 vol.%.
In a L-bending test (r = 0) of the cold-rolled steel
plates, the number of cracks in a corner portion, observed
25 with the naked eye, is preferably two or less per meter.
Hereinafter, production conditions of the cold-rolled
steel plates according to embodiments of the present
disclosure will be described.
Steel slabs having the above-described composition are
30 heated, then hot-finishing rolled at an Ar3 transformation
12
point temperature or above, and coiled at a temperature of
500-800°C.
While the present disclosure particularly limits the
heating temperature of steel slabs, it is preferable that the
heating temperature of steel slabs be limited to 1100°5 C or
above in order to allow the hot rolling-finishing temperature
to be stably obtained.
The hot rolling-finishing temperature is preferably
limited to an Ar3 transformation point temperature or above
10 in order to perform the rolling in an austenite single phase
temperature region.
More preferably, the hot rolling-finishing temperature
is Ar3-950°C.
In the hot-finishing rolling, the reduction ratio and
15 the cooling conditions are not particularly limited. The
coiling temperature is preferably limited to 500°C or above
in order to obtain cold rollability, but the upper limit
thereof is preferably limited to 800°C for preventing grain
coarseness.
20 The thickness of the hot-rolled steel plate is not
particularly limited but is preferably 1.0-3.0 mm.
In the present disclosure, a precipitation
strengthening element is not added in a large amount, and the
coiling temperature is controlled to 500°C or above so that a
25 hard texture is not formed during hot-rolling, the final
strength of hot-rolled steel plates is not excessively high,
and the rolling load of a picking & cold rolling mill (PCM)
during cold rolling is decreased.
Next, the hot-rolled steel plates are cold-rolled at a
30 reduction ratio of 50-90%, and a continuous annealing process,
13
in which the cold-rolled steel plates are maintained in a
continuous annealing line at an annealing temperature of 750-
850°C for 30 seconds or more, these steel plates are then
cooled at a cooling rate of 10-50°C/sec to a temperature
region (an overaging temperature region) within a range 5 of
250-450°C and are maintained at the temperature for 50
seconds or more (overaging), and cooling is finally performed,
thereby producing thin cold-rolled steel plates having high
strength and high formability.
10 The reduction ratio of cold rolling is set to determine
the thickness of a final steel sheet, and is preferably
limited to 50-90%.
When the reduction ratio of cold rolling is below 50%,
it is difficult to obtain a targeted thickness, and when the
15 reduction ratio exceeds 90%, rollability may be reduced.
When the annealing temperature is lower than 750°C,
reverse transformation to austenite may not sufficiently
occur, and when the annealing temperature exceeds 850°C, heat
buckling may easily occur.
20 When the maintenance time is less than 30 seconds,
reverse transformation to austenite may not sufficiently
occur. Therefore, the maintenance time is preferably limited
to 30 seconds or more.
In the case that the cooling stop temperature
25 (overaging temperature) is lower than 250°C or exceeds 450°C,
bainite is not sufficiently formed. Therefore, the cooling
stop temperature (overaging temperature) is preferably
limited to 250-450°C.
When the cooling rate is below 10°C/sec, pearlite may
30 be formed, when the cooling rate exceeds 50°C/sec, martensite
14
may be formed. Therefore, the cooling rate is preferably
limited to 10-50°C/sec.
A more preferable cooling rate is 10-30°C/sec.
When the maintenance time (overaging time) is less than
50 seconds, bainite may not be sufficiently 5 formed.
Therefore, the maintenance time (overaging time) is
preferably limited to 50 seconds or more.
The moving velocity of the steel plates during the
continuous annealing is preferably limited to 100-500 m/min
10 in order to form a fine bainite phase.
In the present invention, a thin cold-rolled steel
plate is produced by cooling a steel material in which
reverse transformation to an austenite phase occurs during
annealing at 750-850°C through control of components
15 contained therein, to a temperature range of 250-450°C in a
state in which the austenite phase is not transformed to a
pearlite phase, or the like during cooling, and maintaining
the steel material within the temperature to cause bainite
transformation, thus forming a low temperature transformation
20 texture in steel.
The cold-rolled steel plates produced as above include
70-100 vol.% of bainite and 0-30 vol.% of ferrite.
In a L-bending test (r = 0) of the cold-rolled steel
plates, the number of cracks in a corner portion, observed
25 with the naked eye, is preferably two or less per meter.
The thickness of the cold-rolled steel plate is not
particularly limited but is preferably 0.5 mm or less.
As described above, according to the present disclosure,
steel plates with targeted strength and formability can be
30 obtained in a continuous annealing line without secondary
15
rolling by using a method in which relatively expensive
elements, such as Mo, Nb, Ti, and the like are excluded,
relatively inexpensive elements, such as Mn, B, and the like
are used to promote bainite transformation during continuous
annealing in a state that an initial strength is 5 not
increased.
Compared with related art inventions in which a rapid
cooling rate of 50°C/sec or above is used in order to cause a
transformation in low carbon steels and a martensite
10 microstructure or the like is formed, steel plates according
to the present disclosure may overcome the low formability
that is the characteristic of such a martensite
microstructure and prevent distortion due to shear
transformation while having a similar level of strength to
15 existing steel plates.
Also, the present disclosure enables low temperature
and high strength transformations to be obtained without the
addition of relatively expensive alloying elements or the
application of rapid cooling rates by decreasing the cooling
20 rate for transformation in a continuous annealing process to
a cooling rate on the continuous annealing line (CAL) level.
Also, since secondary rolling is not performed in the
present disclosure, steel plates according to the present
disclosure have good formability in an L-bending process that
25 is a deformation mode for high strength thin materials, and
high yield strength (YR).
Hereinafter, the present invention will be described in
more detail with examples thereof.
(Example 1)
30 Steels having the compositions shown in Table 1 were
16
hot-rolled (heating temperature: 1250°C, finish-rolling
temperature: 900°C, thickness of hot-rolled steel plates: 2.7
mm, and coiling temperature: 600°C), then these hot-rolled
steel plates were cold-rolled (reduction ratio of primary
cold rolling: 89%, thickness: 0.3 mm) under the 5 production
conditions shown in Table 2 below, then these cold-rolled
steel plates were annealed under the production conditions
shown in Table 3 below to examine yield strength, overall
elongation, hardness, and formability (existence or non10
existence of cracks in L-bending), and yield strength and the
overall elongation are provided in Table 2, hardness is
provided in Table 4, and formability evaluation results
(existence or non-existence of cracks) are provided in Table
5, respectively.
15 Transformation amounts of inventive steels were
measured according to annealing conditions, and the results
are provided in Table 6 below.
In Table 2, yield strength and elongation in
comparative steels A and B are values according to secondary
20 reduction ratios, and yield strength and elongation in
inventive steels are values directly after continuous
annealing without secondary rolling.
Table 5 shows formability test results in inventive
steels and comparative steels in which, since the L-bending
25 test is influenced by whether or not cracks are formed in die
clearances, a poor environment in which an interval between
dies was set to almost zero was assumed and 90 degrees Lbending
text was performed using r=0 bending.
Since the inventive steels made at an annealing
30 temperature of about 700°C failed to obtain a targeted degree
17
of high strength (reverse transformation did not sufficiently
occur due to a low annealing temperature, so that the
fraction of bainite in the microstructure was low), the
annealing temperatures for formability test of test pieces
were limited to 750°C, 780°C, and 800°C, respectively. 5 The
test was performed twice in total. In Table 5, ○ indicates
cracking generation, △ indicates that cracking was not
generated but necking, a precursor to cracking, was generated,
while X indicates a clear surface in which cracking was not
10 generated.
Table 6 shows relative amounts of bainite
transformations in an overaging temperature region of
approximately 350°C obtained by dilatometer measurements in
order to simulate the influence of the annealing conditions
15 of the inventive steels on phase transformation. In this
regard, the last term, i.e., the normalized transformation
length indicates a relative amount of transformation from
austenite to bainite in the overaging temperature region of
approximately 350°C.
20 In Tables 2 and 5 below, B indicates bainite, F
indicates ferrite, and P indicates pearlite.
[Table 1]
Kinds
of
steels
Composition (wt%)
C Mn Si P S Al Ti B N
Comp. A 0.18 0.76 0.012 0.016 0.0044 0.036 0.01 - 0.0034
Comp. B 0.18 1.23 0.012 0.08 0.0049 0.021 - - 0.0042
Inv. 0.19 2.24 0.17 0.01 0.008 0.03 0.016 0.0016 0.0053
*Comp. A: Comparative steel A, Comp. B: Comparative steel B,
25 Inv.: Inventive steel
18
[Table 2]
S.R. (%) 0 6 10 14 20 25 30 40 50 54 Texture
Comp
. A
Y.S.
(MPa)
390 - - 499 - 516 - 636 706 730 F single
phase
Tot.
Elong.
25 - - 6.4
1
- 3.2
2
- 1.5
4
1.8
8
3.1
Comp
. B
Y.S.
(MPa)
447 553 685 657 720 - 770 835 - - F+P
Tot.
Elong.
23.7 14.8 6.0 9.5 5.2 - 3.9 3.6 - -
Inv.
Con.
A
Y.S.
(MPa)
495 - - - - - - - - - B+F(20
vol.%)
Tot.
Elong.
4.34 - - - - - - - - -
Inv.
Con.
B
Y.S.
(MPa)
650 - - - - - - - - - B+F (9
vol.%)
Tot.
Elong.
5.0 - - - - - - - - -
Inv.
Con.
C
Y.S.
(MPa)
791 - - - - - - - - - B single
phase
Tot.
Elong.
6.23 - - - - - - - - -
*S.R.: Secondary reduction ratio, Comp. A: Comparative steel
A, Comp. B: Comparative steel B, Inv. Con. A: Inventive steel
condition A, Inv. Con. B: Inventive steel condition B, Inv.
Con. C: Inventive steel condition C, Y.S.: 5 Yield strength,
Tot. Elong.: Overall elongation
[Table 3]
Inventive
steel
condition
Heating
rate
(°C/sec.)
Cracking
temp.
(°C)
Cracking
Time
(sec.)
Cooling
rate
(°C/sec.)
Overaging
Temp.
(Cooling
step
temp.)
Overaging
time
(maintenance
time) (sec.)
Condition
A
7 700 97 15 350 217
Condition
B
7 750 97 15 350 217
Condition
C
7 800 97 15 350 217
10 [Table 4]
Hardness
(HV500g)
Comp. A
(40%
secondary
reduction
Comp. B
(10%
secondary
reduction
Inv. con.
A
Inv. con.
B
Inv. con.
C
19
) )
Hardness
(HV500g)
220 212 255 367 401
* Comp. A: Comparative steel A, Comp. B: Comparative steel B,
Inv. Con. A: Inventive steel condition A, Inv. Con. B:
Inventive steel condition B, Inv. Con. C: Inventive steel
condition C
5
[Table 5]
Cracking (,
Δ, )
L-bending 180 degree folding Texture
Cooling
rate(°C/sec.)
10 15 20 30 10 15 20 30
Inv.
A.T.:750°C
                B+F
Inv.
A.T.:780°C
    Δ Δ  Δ    Δ    
Inv.
A.T.:800°C
    Δ Δ          
Crack (, Δ,
)
L-bending 180 degree folding Texture
S.R.(%) 10 20 30 40 10 20 30 40
Comp. A                 F
single
phase
Comp. B     Δ Δ     Δ Δ     F+P
*Inv. A.T.: Inventive annealing temperature, S.R.: Secondary
reduction ratio, Comp. A: Comparative steel A, Comp. B:
Comparative steel B
10
[Table 6]
Cracking
(Annealing)
temperature
(°C)
Cracking
(maintenance)
time (sec)
Cooling
rate
(°C/sec)
LCS in BTS
(dL)(mm/mm)
ILS
(LO)
(mm)
Normalized
750 146 15 0.00057 10.2 56.1
750 146 5 0.00053 10.1 52.2
750 291 5 0.00078 10.1 77.3
750 49 10 0.00054 10.4 52.0
750 49 15 0.00052 10.1
2
51.1
750 49 20 0.00072 10.2 71.3
20
750 49 50 0.00147 10.2 143.9
750 97 10 0.00084 10.1
2
82.7
750 97 15 0.00047 10.2 45.8
750 97 20 0.00105 10.2 103.5
750 97 50 0.00197 10.2 193.3
798.5 49 5 0.00108 10.3 105.0
847.5 49 10 0.00235 10.2 230.0
799.5 49 15 0.00317 10.2 310.9
819.5 49 20 0.00321 10.2 314.6
782 49 50 0.00331 10.2 324.8
783.5 97 5 0.00075 10.3 73.0
802 146 5 0.00142 10.2 139.6
*LCS in BTS: Length change of sample in bainite
transformation temperature region, ILS: Initial length of
sample
As provided in Table 2, it can be seen that inve5 ntive
conditions B and C steels other than inventive condition A
steel have the same or higher physical properties such as
combination of yield strength and elongation than comparative
example B after secondary rolling. For example, in order for
10 comparative steel A to obtain a yield strength of 630 MPa or
more, secondary rolling of 40% or more should be performed
and elongation obtained at the second rolling is only 1.5%.
Also, in order for comparative steel B to obtain a
yield strength similar to that of comparative steel A,
15 secondary rolling of 6-10% should be performed thereon, and
elongation obtained within the second rolling range is about
6% and is relatively high.
On the other hand, it can be seen that the steels of
inventive conditions B and C, not subject to secondary
20 rolling and being annealed at a temperature of 750°C or above,
had yield strengths above 650 MPa and elongation above 5.0.
21
Meanwhile, it can be seen that since the steel of
inventive condition A had a low annealing temperature, the
steel of inventive condition A had a lower yield strength
than the steels of inventive conditions B and C and had
5 ductility of 5% or less.
Meanwhile, since real thin steel sheets are extremely
thin and thus may have errors in yield strength, hardness is
frequently used as a standard for the measurement of strength,
instead of yield strength.
10 As provided in Table 4, it can be seen that compared
with yield strength, hardness of inventive steels was
remarkably higher than that of comparative steels.
This phenomenon may be interpreted in the light of
steel hardness generally being proportional not to yield
15 strength, but to tensile strength, and compared with
comparative steels A and B which were subject to secondary
rolling and partially work-hardened through the secondary
rolling, inventive steels which were not subject to secondary
rolling and thus not work-hardened are based on a bainite
20 microstructure having a predetermined strength by itself and
thus show high hardness characteristics.
For example, in tensile tests, the tensile strengths of
comparative steels A and B were not increased by 30 MPa or
more compared with the yield strengths thereof, whereas the
25 tensile strength of the steel of inventive condition A was
683 MPa with respect to the yield strength of 495 MPa, the
tensile strength of the steel of inventive condition B was
949 MPa with respect to the yield strength of 650 MPa, and
the tensile strength of the steel of inventive condition C
30 was 1038 MPa with respect to the yield strength of 790 MPa.
22
That is, all inventive condition steels have remarkably
higher tensile strengths than comparative steels having a
tensile strength of about 700 MPa. It can be understood that
such a high tensile strength guarantees a high degree of
hardness and this effect is much higher in actual 5 ultrathin
steels. These differences in physical properties are due to
a difference in texture between inventive steels and
comparative steels.
The reason why the inventive steels have a greater
10 hardness than the comparative steels is due to the bainite
microstructure included in the inventive steels.
That is, since the comparative steels have a dual phase
structure of ferrite + pearlite and are subject to secondary
rolling, the comparative steels have increased strengths and
15 decreased elongation, whereas since the inventive steels are
not subject to secondary rolling, the inventive steels can
maintain inherent levels of elongation and strength due to
inherent microstructure characteristics thereof. In
conclusion, the inventive steels may have physical properties
20 that are the same as or higher than those of the comparative
steels.
As provided in Table 5, it can be confirmed that in the
inventive steel which was subject to annealing at 750°C,
cracks were generated at various cooling rates and thus the
25 inventive steel was broken, but in the inventive steels which
were subject to an annealing at 780°C or above and were
cooled at a low cooling rate of 15°C/sec, the inventive
steels were not broken in L-bending or even in poorer folding.
On the other hand, in the case of comparative steel A,
30 subject to secondary rolling at a reduction ratio of about
23
40%, cracks were formed after bending and thus comparative
steel A was broken, and in the case of comparative steel B,
subject to secondary rolling at a reduction ratio of 10% or
less, cracks were not formed, but when the reduction ratio
was 5 increased, cracks were formed or necking occurred.
Thus, it is confirmed that the inventive steels which
were subject to annealing under the conditions of an
annealing temperature of 780°C or above and a cooling rate of
15°C/sec have the formability that is equivalent to or higher
10 than that of comparative steel B.
As provided in Table 6, the higher the cooling rate
after annealing at the same annealing temperature and time
annealing, the greater the transformation amount to bainite
is, but since the annealing time which was determined as an
15 important factor to cause reverse transformation did not
almost show an effect, it can be expected that a maintenance
time of 30 seconds or more at a temperature of 750°C or above
allows for sufficient occurrence of reverse transformation
from ferrite to austenite.
20 Meanwhile, it has been found that the influence of the
annealing temperature is very high, and the higher the
annealing temperature, the much more the fraction of
transformation to bainite.
In an aspect of phase transformation, while an
25 annealing temperature of 800°C and a fast cooling rate enable
the active formation of a bainite phase, the continuous
annealing equipment can sufficiently form the bainite phase
even at a low cooling rate of 20°C/sec at the present.
Therefore, the annealing conditions for the inventive steels
30 are limited to an annealing temperature of 750°C or above and
24
an annealing cooling rate of 10-50°C/sec.
Thus, the inventive steels have advantages in that an
additional process such as secondary rolling may be omitted,
the formability in the rolling direction is superior, the
annealing conditions are typical 5 ical continuous annealing
conditions for production of typical products, and high
strength of 900 MPa or more that is the level of tensile
strength may be obtained.
(Example 2)
10 Optical photomicrographs and scanning electron
microscopic (SEM) photographs of inventive steel and
comparative steel B in Example 1 are illustrated in FIGS. 1
and 2, respectively.
In detail, FIG. 1A shows a photomicrograph of inventive
15 steel and FIG. 1B shows a photomicrograph of comparative
steel. FIG. 2 provides scanning electron photomicrographs of
inventive materials taken at magnifications of 1000x (FIG.
2A), 2000x (FIG. 2B), and 5000x (Fig. 2C).
The inventive steel was produced under annealing
20 condition C in which annealing is performed at 800°C, and the
comparative steel B was produced through annealing and then
secondary rolling at a reduction ratio of 14%.
From FIG. 1, a difference in texture between the
inventive steel and the comparative steel can be apparently
25 confirmed. That is, it is illustrated that the comparative
steel has a mixed dual-phase structure of black pearlite and
ferrite, whereas the inventive steel has a single phase
acicular structure.
In order to grasp such textural characteristics, the
30 texture of the inventive steel was observed by a scanning
25
electron microscopy of high magnification and is illustrated
in FIG. 2. From the photographs of FIG. 2 taken at
magnifications of 1000x (FIG. 2A), 2000x (Fig. 2B), and 5000x
(FIG. 2C), it can be confirmed that these inventive steels
have a typical bainite texture in which carbides are 5 formed
inside acicular ferrite laths.
While exemplary embodiments have been shown and
described above, it will be apparent to those skilled in the
art that modifications and variations could be made without
10 departing from the spirit and scope of the present disclosure
as defined by the appended claims.
26
We Claim:
1. A thin cold-rolled steel plate having high
strength and high formability, comprising: 0.15-0.25 wt% of
carbon (C); 1.5-2.5 wt% of manganese (Mn), 0.1-1.0 wt% 5 of
silicon (Si), 0.01-0.05 wt% of titanium (Ti), 5-30 ppm of
boron (B), and a balance of Fe and inevitable impurities,
wherein the ultrathin cold-rolled steel sheet has a
microstructure comprising 70-100 vol.% of bainite, and 0-30
10 vol.% of ferrite.
2. The thin cold-rolled steel plate of claim 1,
wherein the product of contents of C, Mn and B satisfies a
relationship of 1.13*10-4 < wt% C * wt% Mn * wt% B <
15 1.875*10-3.
3. The thin cold-rolled steel plate of claim 1 or 2,
wherein the cold-rolled steel sheet has a thickness less than
0.5 mm.
20
4. The thin cold-rolled steel plate of claim 1 or 2,
wherein in a L-bending test (r = 0) of the cold-rolled steel
sheet, the number of cracks in a corner portion of the coldrolled
steel sheet, observed with the naked eye, is two or
25 less per meter.
5. A method of producing a thin cold-rolled steel
plate, comprising:
heating a steel slab including 0.15-0.25 wt% of carbon
30 (C), 1.5-2.5 wt% of manganese (Mn), 0.1-1.0 wt% of silicon
27
(Si), 0.01-0.05 wt% of titanium (Ti), 5-30 ppm of boron (B),
and a balance of Fe and inevitable impurities;
performing hot-finish rolling of the steel slab at a
temperature equal to an Ar3 transformation point or above to
form a hot-rolled steel plate5 ;
coiling the hot-rolled steel plate at a temperature of
500-800°C;
cold-rolling the hot-rolled steel plate at a reduction
ratio of 50-90% to form a cold-rolled steel plate;
10 maintaining the cold-rolled steel plate in a continuous
annealing line at an annealing temperature of 750-850°C for
30 seconds or more;
cooling the cold-rolled steel plate at a temperature of
250-450°C at a cooling rate of 10-50°C/sec;
15 maintaining the cooled steel plate at the temperature
of 250-450°C for 50 seconds or more; and
recooling the steel plate.
6. The method of claim 5, wherein the product of
20 contents of C, Mn and B satisfies a relationship of 1.13*10-4
< wt% C * wt% Mn * wt% B < 1.875*10-3.
7. The method of claim 5 or 6, wherein the steel
plate during the continuous annealing has a moving velocity
25 at 100-500 m/min.
8. The method of claim 5 or 6, wherein the hotfinish
rolling temperature is Ar3-950°C and the cooling rate
is 10-30°C/sec.
30
28
9. The method of claim 5 or 6, wherein the hotrolled
steel plate has a thickness in a range of 1.0-3.0 mm
and the cold-rolled steel plate has a thickness less than 0.5
mm.
5
10. The method of claim 5 or 6, wherein in a Lbending
test (r = 0) of the cold-rolled steel plate, the
number of cracks in a corner portion of the cold-rolled steel
plate, observed with the naked eye, is two or less per meter.