U.S. patent application number 15/565096 was filed with the patent office on 2018-04-19 for steel sheet with excellent cold workability during forming and method for manufacturing the same.
This patent application is currently assigned to NIPPON STEEL & SUMITOMO METAL CORPORATION. The applicant listed for this patent is NIPPON STEEL & SUMITOMO METAL CORPORATION. Invention is credited to Motonori HASHIMOTO, Kazuo HIKIDA, Ken TAKATA, Kengo TAKEDA.
Application Number | 20180105891 15/565096 |
Document ID | / |
Family ID | 57072674 |
Filed Date | 2018-04-19 |
United States Patent
Application |
20180105891 |
Kind Code |
A1 |
HIKIDA; Kazuo ; et
al. |
April 19, 2018 |
STEEL SHEET WITH EXCELLENT COLD WORKABILITY DURING FORMING AND
METHOD FOR MANUFACTURING THE SAME
Abstract
The present invention provides a steel sheet having an excellent
cold workability during forming and a method for producing the
same. The steel sheet of the present invention is characterized in
that: (a) the ratio of the number of carbides at the ferrite grain
boundary to the number of carbides in the ferrite grain exceeds 1,
(b) the ferrite grain diameter is 5 .mu.m or more and 50 .mu.m or
less, (c) the in-plane anisotropy |.DELTA.r| of the r value is 0.2
or less, (d) the Vickers hardness is 100 HV or more and 150 HV or
less, (e) the random intensity ratio of the {311} <011>
orientation at the 1/2-thickness portion of the steel sheet is 3.0
or less.
Inventors: |
HIKIDA; Kazuo; (Tokyo,
JP) ; HASHIMOTO; Motonori; (Tokyo, JP) ;
TAKEDA; Kengo; (Tokyo, JP) ; TAKATA; Ken;
(Tokyo, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
NIPPON STEEL & SUMITOMO METAL CORPORATION |
Tokyo |
|
JP |
|
|
Assignee: |
NIPPON STEEL & SUMITOMO METAL
CORPORATION
Tokyo
JP
|
Family ID: |
57072674 |
Appl. No.: |
15/565096 |
Filed: |
April 8, 2016 |
PCT Filed: |
April 8, 2016 |
PCT NO: |
PCT/JP2016/061608 |
371 Date: |
October 6, 2017 |
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C21D 2211/005 20130101;
C22C 38/06 20130101; C22C 38/001 20130101; C22C 38/14 20130101;
C22C 38/10 20130101; C22C 38/002 20130101; C22C 38/18 20130101;
C22C 38/12 20130101; C22C 38/02 20130101; C22C 38/60 20130101; C21D
8/0226 20130101; C21D 9/46 20130101; C22C 38/04 20130101; C22C
38/16 20130101; C21D 8/0205 20130101; C21D 8/0263 20130101; C22C
38/005 20130101; C22C 38/08 20130101 |
International
Class: |
C21D 8/02 20060101
C21D008/02; C21D 9/46 20060101 C21D009/46; C22C 38/06 20060101
C22C038/06; C22C 38/04 20060101 C22C038/04; C22C 38/02 20060101
C22C038/02 |
Foreign Application Data
Date |
Code |
Application Number |
Apr 10, 2015 |
JP |
2015-081101 |
Claims
1. A steel sheet having an excellent cold workability during
forming, comprising, in terms of % by mass: C: 0.10 to 0.40%, Si:
0.01 to 0.30%, Mn: 0.30 to 1.00%, P: 0.0001 to 0.020%, S: 0.0001 to
0.010%, Al: 0.001 to 0.10%, and a balance of Fe and inevitable
impurities, wherein (a) a ratio of the number of carbides at a
ferrite grain boundary relative to the number of carbides in the
ferrite grain is more than 1, wherein (b) a diameter of the ferrite
grain is 5 .mu.m or more and 50 .mu.m or less, wherein (c) an
in-plane anisotropy |.DELTA.r| of the r value standardized
according to JIS Z 2254 is 0.2 or less, wherein (d) a Vickers
hardness of the steel sheet is 100 HV or more and 150 HV or less,
and wherein (e) a ratio of X-ray diffraction intensity of the {311}
<011> orientation at the 1/2-thickness portion of the steel
sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in
the steel sheet is subjected to X-ray diffraction is 3.0 or
less.
2. The steel sheet with excellent cold workability during forming
according to claim 1 further comprising, in terms of % by mass, one
or a plurality of: N: 0.0001 to 0.010%, O: 0.0001 to 0.020%, Cr:
0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to 0.10%, V: 0.001 to
0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001 to 0.10%,
Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As:
0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y:
0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, and Ce:
0.001 to 0.050%.
3. A method for producing a steel sheet with excellent cold
workability during forming according to claim 1, said method
comprising: subjecting a steel strip having an ingredient
composition according to claim 1 to hot rolling by heating,
followed by completing the finish hot rolling at a temperature
range of 800.degree. C. or higher and 900.degree. C. or lower;
coiling said hot-rolled steel sheet at a temperature of 400.degree.
C. or higher and 550.degree. C. or lower; pickling said hot-rolled
steel sheet, and then subjecting said hot-rolled steel sheet to a
two-step type annealing in which said hot-rolled steel sheet is
retained in two temperature ranges, wherein the two-step type
annealing comprises (i) subjecting said hot-rolled steel sheet to a
first step annealing performed by retaining said hot-rolled steel
at a temperature range of 650.degree. C. or higher and 720.degree.
C. or lower for 3 hours or longer and 60 hours or shorter, and then
a second step annealing performed by retaining the hot-rolled steel
at a temperature range of 725.degree. C. or higher and 790.degree.
C. or lower for 3 hours or longer and 50 hours or shorter, and
thereafter (ii) cooling said hot-rolled steel sheet to 650.degree.
C. or lower at a cooling rate of 1.degree. C./hour or more and
30.degree. C./hour or less.
4. The method for producing a steel sheet according to claim 3,
wherein the steel sheet has a cross-sectional shrinkage percentage
of 40% or more.
5. A method for producing a steel sheet with excellent cold
workability during forming according to claim 2, said method
comprising: subjecting a steel strip having an ingredient
composition according to claim 2 to hot rolling by heating,
followed by completing the finish hot rolling at a temperature
range of 800.degree. C. or higher and 900.degree. C. or lower;
coiling said hot-rolled steel sheet at a temperature of 400.degree.
C. or higher and 550.degree. C. or lower; pickling said hot-rolled
steel sheet, and then subjecting said hot-rolled steel sheet to a
two-step type annealing in which said hot-rolled steel sheet is
retained in two temperature ranges, wherein the two-step type
annealing comprises (i) subjecting said hot-rolled steel sheet to a
first step annealing performed by retaining said hot-rolled steel
at a temperature range of 650.degree. C. or higher and 720.degree.
C. or lower for 3 hours or longer and 60 hours or shorter, and then
a second step annealing performed by retaining the hot-rolled steel
at a temperature range of 725.degree. C. or higher and 790.degree.
C. or lower for 3 hours or longer and 50 hours or shorter, and
thereafter (ii) cooling said hot-rolled steel sheet to 650.degree.
C. or lower at a cooling rate of 1.degree. C./hour or more and
30.degree. C./hour or less.
6. The method for producing a steel sheet according to claim 5,
wherein the steel sheet has a cross-sectional shrinkage percentage
of 40% or more.
Description
TECHNICAL FIELD
[0001] The present invention relates to a steel sheet with
excellent cold workability during forming and a method for
manufacturing the sheet.
BACKGROUND ART
[0002] Automotive parts, knives, and other mechanical parts are
manufactured through working processes such as punching, bending,
and pressing. In the working processes, improvement of workability
is required for a material carbon steel sheet, in order to improve
product quality and stability and/or cost reduction.
[0003] Generally, a carbon steel sheet is subjected to cold rolling
and spheroidizing annealing, so as to produce a soft carbon steel
sheet with excellent workability made of ferrite and spheroidized
carbide. Many technologies for improving the workability of carbon
steel sheets have been proposed so far.
[0004] For example, Patent Document 1 discloses a high-carbon steel
sheet for precision punching and a method for producing the sheet,
wherein the sheet comprises, in terms of % by mass, C: 0.15 to
0.90%, Si: 0.40% or less, Mn: 0.3 to 1.0%, P: 0.03% or less, total
Al: 0.1% or less, Ti: 0.01 to 0.05%, B: 0.0005 to 0.0050%, N: 0.01%
or less, and Cr: 1.2% or less, has a structure in which carbides
having an average carbide grain size of 0.4 to 1.0 .mu.m and a
carbide spheroidization ratio of 80% or more are dispersed in a
ferrite matrix, and has a notched tensile elongation of 20% or
more.
[0005] Patent Document 2 discloses a medium- to high-carbon steel
sheet with excellent workability and a method for producing the
sheet, wherein the sheet comprises C: 0.3 to 1.3 wt %, Si: 1.0 wt %
or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt% or less, and S: 0.02 wt %
or less, has a structure in which carbides are dispersed so that
the relationship C.sub.GB/C.sub.IG.ltoreq.0.8 holds between the
carbide number C.sub.GB on the ferrite crystal grain boundary and
the carbide number C.sub.IG in the ferrite crystal grains, and has
a cross-sectional hardness of 160 HV or less.
[0006] Patent Document 3 discloses a medium- to high-carbon steel
sheet with excellent workability, wherein the sheet comprises C:
0.30 to 1.00 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P:
0.02 wt % or less, and S: 0.02 wt % or less, has a structure in
which carbides are dispersed in ferrite so that the relationship
C.sub.GB/C.sub.IG.ltoreq.0.8 holds between the carbide number
C.sub.GB on the ferrite crystal grain boundary and the carbide
number C.sub.IG in the ferrite crystal grains, and simultaneously
90% or more of the total carbides are occupied by spheroidized
carbides having a long axis/short axis of 2 or less.
[0007] Patent Documents 1 to 3 describe that the greater the
proportion of carbides in ferrite grains, the more the workability
is improved.
[0008] In addition, Patent Document 4 discloses a steel sheet
having excellent FB workability, mold life, and cold formability
after FB processing, wherein the sheet comprises C: 0.1 to 0.5 wt
%, Si: 0.5 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.03 wt % or less,
S: 0.02 wt % or less, has a structure based on ferrite and carbide,
and the amount S.sub.gb of the carbide present on the ferrite grain
boundary is 40% or more, the above S.sub.gb being defined by
S.sub.gb={S.sub.on/(S.sub.on+S.sub.in)}.times.100 (wherein S.sub.on
is the total area occupied by the carbides present on the grain
boundary among the carbides present per unit area and S.sub.in is
the total area occupied by the carbides present on the grain
boundary among the carbides present per unit area).
[0009] However, in the technology described in Patent Document 1,
annealing is performed at a temperature of the A.sub.C1 point or
higher for softening in order to coarsen ferrite grain size and
carbide. But when annealing is performed at a temperature of the
A.sub.C1 point or higher, rod-like/plate-like carbides may
precipitate during annealing. The carbides, even though capable of
reducing hardness, deteriorate workability, which is
disadvantageous in terms of workability.
[0010] The technologies described in Patent Documents 2 and 3
consider that the deterioration of workability is caused by the low
carbide spheroidization ratio of carbides precipitated on the grain
boundary, but do not take into account the problem of improving the
spheroidization ratio of grain boundary carbides. Techniques
described in Patent Document 4 only specify the tissue factor, and
Patent Document 4 does not discuss the relationship between
workability and mechanical properties.
[0011] The technology described in Patent Document 5 is an
invention made by focusing on the relationship between fine
blanking workability and the amount of carbide present in ferrite
grains and ferrite grain size. However, Patent Document 5 does not
discuss what effect the aggregate structure has on the plastic
anisotropy.
[0012] Patent Document 6 discloses a hot-rolled steel sheet in
which the development of an aggregate structure otherwise developed
by rolling is suppressed and a method for manufacturing the sheet.
However, Patent Document 6 does not discuss the relationship
between the aggregate structure other than the aggregate structure
developed by rolling and the cold forgeability.
[0013] The technology described in Patent Document 7 is an
invention made by considering that the hardness and the total
elongation of a high-carbon hot-rolled steel sheet prior to
quenching are greatly influenced by the cementite density in the
ferrite grains. The hot-rolled steel sheet described in Patent
Document 7 is characterized in that it has a microstructure
composed of ferrite and cementite, said microstructure having a
cementite density of 0.10 strips/.mu.m.sup.2 or less in the ferrite
grains. However, Patent Document 7 does not discuss what effect the
aggregate texture has on the plastic anisotropy.
[0014] The technology described in Patent Document 8 is an
invention made by considering that the C.sub.eq value is related
not only to mechanical properties and weldability but also to the
fatigue crack growth rate in steels having a fine structure. Patent
Document 8 discloses that by limiting the range of the C.sub.eq
value to a range of 0.28% to 0.65%, the fatigue resistance of the
steel material is improved and simultaneously weldability is
secured. However, Patent Document 8 does not discuss what effect
the aggregate texture has on the plastic anisotropy.
PRIOR ART DOCUMENTS
Patent documents
[0015] [Patent Document 1] Japanese Patent No. 4465057
[0016] [Patent Document 2] Japanese Patent No. 4974285
[0017] [Patent Document 3] Japanese Patent No. 5197076
[0018] [Patent Document 4] Japanese Patent No. 5194454
[0019] [Patent Document 5] Japanese Unexamined Patent Publication
No. 2007-270331
[0020] [Patent Document 6] Japanese Unexamined Patent Publication
No. 2009-263718
[0021] [Patent Document 7] Japanese Unexamined Patent Publication
No. 2015-17294
[0022] [Patent Document 8] Japanese Unexamined Patent Publication
No. 2004-27355
SUMMARY OF THE INVENTION
Problem to be Solved by the Invention
[0023] In view of the current state of the prior art, it is an
object of the present invention to address the problem of improving
the cold workability of a steel sheet during forming, and to
provide a steel sheet that has solved the problem and a method for
manufacturing the sheet.
Means to Solve the Problem
[0024] The present inventors have conducted intensive and extensive
studies on methods for solving the above-mentioned problems. As a
result, the present inventors have found that by controlling the
dispersion state of the carbide in the structure of the steel sheet
before cold working through the optimization of the manufacturing
conditions in the steps from hot rolling to annealing, the carbide
can be precipitated on the ferrite boundary and simultaneously the
aggregate structure in the hot rolled steel plate can be
controlled, thereby leading to enhanced cold workability.
[0025] Further, we have found after intensive and extensive
research that it is difficult to manufacture a steel sheet that
satisfies the above-mentioned conditions merely by devising hot
rolling conditions and annealing conditions separately, and that it
can be manufactured by optimizing the above conditions in mutual
cooperation in an integrated process of the hot rolling and
annealing steps.
[0026] The present invention has been made based on the above
findings, and the gist thereof lies in:
[0027] (1) A steel sheet having an excellent cold workability
during forming, comprising, in terms of % by mass:
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.0001 to 0.010%,
Al: 0.001 to 0.10%, and
[0028] a balance of Fe and inevitable impurities, wherein (a) a
ratio of the number of carbides at a ferrite grain boundary
relative to the number of carbides in the ferrite grain is more
than 1, wherein (b) a diameter of the ferrite grain is 5 .mu.m or
more and 50 .mu.m or less, wherein (c) an in-plane anisotropy
|.DELTA.r| of the r value standardized according to JIS Z 2254 is
0.2 or less, wherein (d) a Vickers hardness of the steel sheet is
100 HV or more and 150 HV or less, and wherein (e) a ratio of X-ray
diffraction intensity of the {311} <011> orientation at the
1/2-thickness portion of the steel sheet relative to the X-ray
diffraction intensity obtained when a sample with a random
orientation distribution of crystal grains in the steel sheet is
subjected to X-ray diffraction is 3.0 or less.
[0029] (2) The steel sheet with excellent cold workability during
forming described in the above (1) further comprising, in terms of
% by mass, one or a plurality of:
N: 0.0001 to 0.010%,
O: 0.0001 to 0.020%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.10%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%,
Cu: 0.001 to 0.10%,
W: 0.001 to 0.10%,
Ta: 0.001 to 0.10%,
Ni: 0.001 to 0.10%,
Sn: 0.001 to 0.050%,
Sb: 0.001 to 0.050%,
As: 0.001 to 0.050%,
Mg: 0.0001 to 0.050%,
Ca: 0.001 to 0.050%,
Y: 0.001 to 0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%, and
Ce: 0.001 to 0.050%.
[0030] (3) A method for producing a steel sheet with excellent cold
workability during forming according to the above (1) or (2), the
method comprising:
[0031] subjecting a steel strip having an ingredient composition
according to claim 1 or 2 to hot rolling by heating, followed by
completing the finish hot rolling at a temperature range of
800.degree. C. or higher and 900.degree. C. or lower;
[0032] coiling the hot-rolled steel sheet at a temperature of
400.degree. C. or higher and 550.degree. C. or lower;
[0033] pickling the hot-rolled steel sheet, and then subjecting the
hot-rolled steel sheet to a two-step type annealing in which the
hot-rolled steel sheet is retained in two temperature ranges,
[0034] wherein the two-step type annealing comprises
[0035] (i) subjecting the hot-rolled steel sheet to a first step
annealing performed by retaining said hot-rolled steel at a
temperature range of 650.degree. C. or higher and 720.degree. C. or
lower for 3 hours or longer and 60 hours or shorter, and then a
second step annealing performed by retaining the hot-rolled steel
at a temperature range of 725.degree. C. or higher and 790.degree.
C. or lower for 3 hours or longer and 50 hours or shorter, and
thereafter
[0036] (ii) cooling the hot-rolled steel sheet to 650.degree. C. or
lower at a cooling rate of 1.degree. C./hour or more and 30.degree.
C./hour or less.
[0037] (4) The method for producing a steel sheet described in the
above (3), wherein the steel sheet has a cross-sectional shrinkage
percentage of 40% or more.
Effect of the Invention
[0038] According to the present invention, a steel sheet with
excellent cold workability during forming can be manufactured and
provided.
Mode for Carrying Out the Invention
[0039] A steel sheet with excellent cold workability during forming
according to the present invention (hereinafter may be referred to
as "the inventive steel sheet") comprises, in terms of % by
mass:
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.0001 to 0.010%,
Al: 0.001 to 0.10%, and
[0040] a balance of Fe and inevitable impurities,
[0041] the above sheet being characterized in that:
(a) the ratio of the number of carbides at a ferrite grain boundary
relative to the number of carbides in the ferrite grain exceeds 1,
(b) the ferrite grain diameter is 5 .mu.m or more and 50 .mu.m or
less, (c) the in-plane anisotropy |.DELTA.r| of the r value
standardized according to JIS Z 2254 is 0.2 or less, (d) the
Vickers hardness is 100 HV or more and 150 HV or less, and (e) the
ratio of X-ray diffraction intensity of the {311}
<011>orientation at the 1/2-thickness portion of the steel
sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in
the steel sheet is subjected to X-ray diffraction is 3.0 or
less.
[0042] The method (hereinafter may be referred to as "the inventive
method") of the present invention for producing a steel sheet with
excellent cold workability during forming is a method for producing
the inventive steel sheet,
[0043] wherein a hot-rolled steel strip that has been obtained by
subjecting a steel strip having an ingredient composition of the
inventive steel sheet to hot rolling by heating, followed by
completing the finish hot rolling at a temperature range of
800.degree. C. or higher and 900.degree. C. or lower, and by
coiling the resulting hot-rolled steel sheet at a temperature of
400.degree. C. or higher and 550.degree. C. or lower is, after
pickling, subjected to two-step type annealing in which the sheet
is retained in two temperature ranges, whereupon
[0044] (i) the hot-rolled steel sheet is subjected to a first step
annealing performed by retaining said hot-rolled steel at a
temperature range of 650.degree. C. or higher and 720.degree. C. or
lower for 3 hours or longer and 60 hours or shorter, and then
subjected to a second step annealing performed by retaining the
hot-rolled steel at a temperature range of 725.degree. C. or higher
and 790.degree. C. or lower for 3 hours or longer and 50 hours or
shorter, and thereafter
[0045] (ii) the sheet is cooled down to 650.degree. C. or lower at
a cooling rate of 1.degree. C./hour or more and 30.degree. C./hour
or less.
[0046] Hereinafter, the inventive steel sheet and the inventive
manufacturing method will be described.
[0047] First, the reasons for limiting the ingredient composition
of the inventive steel sheet will be described. The percentage
relating to the ingredient composition means % by mass.
C: 0.10 to 0.40%
[0048] C is an element that forms carbide in steel, and is
effective for strengthening steel and refining ferrite grains. In
order to prevent the surface of the steel sheet from being textured
by cold working and ensure the aesthetic appearance of surface of
cold forged parts, it is necessary to suppress the coarsening of
ferrite grain size. However, when its content is less than 0.10%,
the volume fraction of the carbide is insufficient and the
coarsening of carbides during annealing cannot be suppressed.
Therefore, C is set to 0.10% or more, and preferably 0.12% or
more.
[0049] On the other hand, when it exceeds 0.40%, the volume
fraction of the carbide increases, a large amount of cracks serving
as fracture starting points are formed when a load is
instantaneously applied, and thus the impact resistance property
decreases. Therefore, C is set to 0.40% or less, and preferably
0.38% or less.
Si: 0.01 to 0.30%
[0050] Si is an element that acts as a deoxidizing agent and also
affects the form of the carbide. In order to reduce the number of
carbides in the ferrite grain and increase the number of carbides
on the ferrite grain boundaries, it is necessary to generate an
austenite phase during annealing in the two-step type annealing,
and, after transiently dissolving the carbides, to cool gradually
to promote the precipitation of carbides at the ferrite grain
boundaries.
[0051] In the inventive steel sheet, the amount of Si may
preferably be as small as possible. However, when it is reduced to
less than 0.01%, the manufacturing cost increases. Therefore, Si is
set to 0.01% or more.
[0052] On the other hand, when it exceeds 0.30%, the ductility of
ferrite lowers and breaking may easily occur during cold working,
resulting in reduced cold workability. Therefore, Si is set to
0.30% or less, and preferably 0.28% or less.
Mn: 0.30 to 1.00%
[0053] Mn is an element that controls the figuration of carbides in
the two-step type annealing. When its content is less than 0.30%,
it is difficult to precipitate carbides at the ferrite grain
boundaries in slow cooling after the second-step annealing.
Therefore, Mn is set to 0.30% or more, and preferably 0.33% or
more.
[0054] On the other hand, when it exceeds 1.00%, the hardness of
ferrite increases and the cold workability deteriorates. Therefore,
Mn is set to 1.00% or less, and preferably 0.96% or less.
P: 0.0001 to 0.020%
[0055] P is an element that segregates at the ferrite grain
boundaries and suppresses the formation of grain boundary carbides.
The amount of P may preferably be as small as possible. However,
when P is reduced to less than 0.0001% in the refining process, the
refining cost may greatly increase. Therefore, it is set to 0.0001%
or more, and preferably 0.0013% or more.
[0056] On the other hand, when it exceeds 0.020%, the number
percentage of the grain boundary carbides decreases and the cold
workability deteriorates. Therefore, P is set to 0.020% or less,
and preferably 0.018% or less.
S: 0.0001 to 0.010%
[0057] S is an element that forms a non-metallic inclusion such as
MnS. Since a non-metallic inclusion serves as the starting point
for break generation during cold forging, the amount of S may
preferably be as small as possible. However, when it is reduced to
less than 0.0001%, the refining cost greatly increases. Therefore,
S is set to 0.0001% or more, and preferably 0.0012% or more.
[0058] On the other hand, when it exceeds 0.010%, cold workability
deteriorates. Therefore, S is set to 0.010% or less, and preferably
0.007% or less.
Al: 0.001 to 0.10%
[0059] Al is an element that acts as a deoxidizing agent for steel
and stabilizes ferrite. When its content is less than 0.001%, a
sufficient addition effect cannot be obtained. Therefore, Al is set
to 0.001% or more, and preferably 0.004% or more.
[0060] On the other hand, when it exceeds 0.10%, the number
percentage of carbides on the grain boundary decreases and the cold
workability deteriorates. Therefore, Al is set to 0.10% or less,
and preferably 0.08% or less.
[0061] In addition to the above elements, the inventive steel sheet
may contain one or a plurality of N: 0.0001 to 0.010%, 0: 0.0001 to
0.020%, Cr: 0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001
to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to
0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to
0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to
0.050%, and Ce: 0.001 to 0.050%, in order to improve the properties
of the inventive steel sheet.
N: 0.0001 to 0.010%
[0062] N is an element that, when present in large amounts, causes
the embrittlement of ferrite. The amount of N may preferably be as
small as possible. However, when it is reduced to less than
0.0001%, the refining cost greatly increases. Therefore, N should
be 0.0001% or more, and preferably 0.0006% or more. On the other
hand, when it exceeds 0.010%, ferrite embrittles and the cold
forgeability deteriorates. Therefore, N should be 0.010% or less,
and preferably 0.007% or less.
O: 0.0001 to 0.020%
[0063] O is an element that, when present in large amounts, forms
coarse oxides in steel. The amount of O may preferably be as small
as possible. However, when it is reduced to less than 0.0001%, the
refining cost increases greatly. Therefore, O is set to 0.0001% or
more, and preferably 0.0011% or more. On the other hand, when it
exceeds 0.020%, coarse oxides are formed in the steel, the oxides
serving as the starting point for break generation during cold
working. Therefore, 0 is set to 0.020% or less, and preferably
0.017% or less.
Cr: 0.001 to 0.50%
[0064] Cr is an element which enhances quenchability and
contributes to the improvement of strength and which is thickened
to carbide and forms stable carbide even in the austenitic phase.
When its content is less than 0.001%, the sufficient effect of
improving quenchability cannot be obtained. Therefore, Cr is set to
0.001% or more, and preferably 0.007% or more. On the other hand,
when it exceeds 0.50%, the carbide becomes stabilized thereby
delaying the dissolution of the carbide during quenching, and thus,
it is feared that the desired quenching strength may not be
achieved. Therefore, Cr is set to 0.50% or less, and preferably
0.45% or less.
Mo: 0.001 to 0.10%
[0065] Like Mn, Mo is an element effective for controlling the
figuration of carbides. When its content is less than 0.001%, a
sufficient addition effect cannot be obtained. Therefore, Mo is set
to 0.001% or more, and preferably 0.010% or more. On the other
hand, when it exceeds 0.10%, the in-plane anisotropy of the r value
deteriorates and the cold workability deteriorates. Therefore, Mo
is set to 0.10% or less, and preferably 0.08% or less.
Nb: 0.001 to 0.10%
[0066] Nb is an element which is effective for controlling the
figuration of carbides and which refines the structure, thereby
contributing to the enhancement of its toughness. When its content
is less than 0.001%, a sufficient addition effect cannot be
obtained. Therefore, Nb should be 0.001% or more, and preferably
0.004% or more. On the other hand, when it exceeds 0.10%, a large
number of fine Nb carbides precipitate, which leads to excessively
increased strength. It also causes the reduction in the number
ratio of grain boundary carbides, and the deterioration in cold
forgeability. Therefore, Nb is set to 0.10 or less, and preferably
0.08% or less.
V: 0.001 to 0.10%
[0067] Like Nb, V is an element which is effective for controlling
the figuration of carbides and which refines the structure, thereby
contributing to the enhancement of its toughness. When its content
is less than 0.001%, a sufficient addition effect cannot be
obtained. Therefore, V is set to 0.001% or more, and preferably
0.004% or more. On the other hand, when it exceeds 0.10%, a large
number of fine V carbides precipitate, which leads to excessively
increased strength, to the reduced number ratio of grain boundary
carbides, and to the deteriorated cold forgeability. Therefore, V
is set to 0.10 or less, and preferably 0.08% or less.
Cu: 0.001 to 0.10%
[0068] Cu is an element which segregates at the ferrite crystal
grain boundary and forms fine precipitates thereby to contribute to
the enhancement of strength. When its content is less than 0.001%,
a sufficient effect of enhancing strength cannot be obtained.
Therefore, Cu is set to 0.001% or more, and preferably 0.005% or
more. On the other hand, when it exceeds 0.10%, red heat
embrittlement occurs and the productivity by hot rolling decreases.
Therefore, Cu is set to 0.10% or less, and preferably 0.08% or
less.
W: 0.001 to 0.10%
[0069] Like Nb and V, W is also an element effective for
controlling the figuration of carbides. When its content is less
than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, W is set to 0.001% or more, and preferably 0.003% or
more. On the other hand, when it exceeds 0.10%, a large number of
fine W carbides precipitate, which leads to excessively increased
strength, to the reduced number ratio of grain boundary carbides,
and to the deteriorated cold forgeability. Therefore, W is set to
0.10 or less, and preferably 0.08% or less.
Ta: 0.001 to 0.10%
[0070] Like Nb, V and W, Ta is also an element effective for
controlling the figuration of carbides. When its content is less
than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, W is set to 0.001% or more, and preferably 0.005% or
more. On the other hand, when it exceeds 0.10%, a large number of
fine W carbides precipitate, which leads to excessively increased
strength, to the reduced number ratio of grain boundary carbides,
and to the deteriorated cold forgeability. Therefore, Ta is set to
0.10 or less, and preferably 0.08% or less.
Ni: 0.001 to 0.10%
[0071] Ni is an element effective for improving the toughness of
parts. When its content is less than 0.001%, a sufficient addition
effect cannot be obtained. Therefore, Ni is set to 0.001% or more,
and preferably 0.003% or more. On the other hand, when it exceeds
0.10%, the number ratio of grain boundary carbides decreases and
the cold forgeability deteriorates. Therefore, Ni is set to 0.10%
or less, and preferably 0.08% or less.
Sn: 0.001 to 0.050%
[0072] Sn is an element contaminated from a steel raw material
(scrap). It segregates at the grain boundary, leading to the
decreased number ratio of grain boundary carbides. Therefore, its
content may preferably be as small as possible. However, when it is
reduced to less than 0.001%, the refining cost will be greatly
increased. Therefore, Sn is set to 0.001% or more, and preferably
0.002% or more. On the other hand, when it exceeds 0.050%, ferrite
embrittles and cold forgeability deteriorates. Therefore, Sn is set
to 0.050% or less, and preferably 0.040% or less.
Sb: 0.001 to 0.050%
[0073] Like Sb, Sb is an element contaminated from a steel raw
material (scrap). It segregates at the grain boundary, leading to
the decreased number ratio of grain boundary carbides. Therefore,
its content may preferably be as small as possible. However, when
it is reduced to less than 0.001%, the refining cost will be
greatly increased. Therefore, Sb is set to 0.001% or more,
preferably 0.002% or more. On the other hand, when it exceeds
0.050%, the cold forgeability deteriorates. Therefore, Sb is set to
0.050% or less, and preferably 0.040% or less.
As: 0.001 to 0.050%
[0074] Like Sn and Sb, As is an element contaminated from a steel
raw material (scrap). It segregates at the grain boundary, thereby
leading to a decrease in the number ratio of grain boundary
carbides. Therefore, its content may preferably be as small as
possible. However, when it is reduced to less than 0.001%, the
refining cost increases greatly. Therefore, As is set to 0.001% or
more, and preferably 0.002% or more. On the other hand, when it
exceeds 0.050%, the number ratio of the grain boundary carbides
decreases and the cold forgeability deteriorates. Therefore, As is
set to 0.050% or less, and preferably 0.040% or less.
Mg: 0.0001 to 0.050%
[0075] Mg is an element that can control the figuration of sulfides
with the addition of its trace amount. When its content is less
than 0.0001%, a sufficient addition effect cannot be obtained.
Therefore, Mg is set to 0.0001% or more, and preferably 0.0008% or
more. On the other hand, when it exceeds 0.050%, ferrite embrittles
and the cold forgeability deteriorates. Therefore, Mg is set to
0.050% or less, and preferably 0.040% or less.
Ca: 0.001 to 0.050%
[0076] Like Mg, Ca is an element that can control the figuration of
sulfides with the addition of its trace amount. When its content is
less than 0.001%, a sufficient addition effect cannot be obtained.
Therefore, Ca is set to 0.001% or more, and preferably 0.003% or
more. On the other hand, when it exceeds 0.050%, coarse Ca oxides
are formed, which serve as starting points of break generation
during cold forging. Therefore, Ca is set to 0.050% or less, and
preferably 0.040% or less.
Y: 0.001 to 0.050%
[0077] Like Mg and Ca, Y is an element that can control the
figuration of sulfides with the addition of its trace amount. When
its content is less than 0.001%, a sufficient addition effect
cannot be obtained. Therefore, Y is set to 0.001% or more, and
preferably 0.003% or more. On the other hand, when it exceeds
0.050%, coarse Y oxides are formed, which serve as starting points
of break generation during cold working. Therefore, Y is set to
0.050% or less, and preferably 0.035% or less.
Zr: 0.001 to 0.050%
[0078] Like Mg, Ca and Y, Zr is an element that can control the
figuration of sulfides with the addition of its trace amount. When
its content is less than 0.001%, a sufficient addition effect
cannot be obtained. Therefore, Zr is set to 0.001% or more, and
preferably 0.004% or more. On the other hand, when it exceeds
0.050%, coarse Zr oxides are formed, which serve as starting points
for break generation during cold working. Therefore, Zr is set to
0.050% or less, and preferably 0.045% or less.
La: 0.001 to 0.050%
[0079] La is an element that can control the figuration of sulfides
with the addition of its trace amount, but it is also an element
that segregates at the grain boundary and causes a decrease in the
number ratio of grain boundary carbides. When its content is less
than 0.001%, a sufficient effect of controlling figuration cannot
be obtained. Therefore, La is set to 0.001% or more, and preferably
0.004% or more. On the other hand, when it exceeds 0.050%, the
number ratio of grain boundary carbides decreases and the cold
workability deteriorates. Therefore, La is set to 0.050% or less,
and preferably 0.045% or less.
Ce: 0.001 to 0.050%
[0080] Like La, Ce is an element that can control the figuration of
sulfides with the addition of its trace amount, but it is also an
element that segregates at the grain boundary and causes a decrease
in the number ratio of grain boundary carbides. When its content is
less than 0.001%, a sufficient effect of controlling figuration
cannot be obtained. Therefore, Ce is set to 0.001% or more, and
preferably 0.004% or more. On the other hand, when it exceeds
0.050%, the number ratio of grain boundary carbides decreases and
the cold forgeability deteriorates. Therefore, Ce is set to 0.050%
or less, and preferably 0.045% or less.
[0081] The remainder of the ingredient composition of the inventive
steel sheet is Fe and unavoidable impurities.
[0082] It is a novel finding by the inventors that the inventive
steel sheet has excellent cold workability during forming, because,
in addition to the above ingredient composition, it was found, as a
result of optimum hot rolling and annealing, that
[0083] (a) the ratio of the number of carbides at the ferrite grain
boundary relative to the number of carbides in the ferrite grain
exceeds 1,
[0084] (b) the ferrite grain diameter is 5 .mu.m or more and 50
.mu.m or less,
[0085] (c) the in-plane anisotropy |.DELTA.r| of the r value
standardized according to JIS Z 2254 is 0.2 or less,
[0086] (d) the Vickers hardness is 100 HV or more and 150 HV or
less, and
[0087] (e) the ratio of X-ray diffraction intensity of the {311}
<011> orientation at the 1/2-thickness portion of the steel
sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in
the steel sheet is subjected to X-ray diffraction is 3.0 or
less.
[0088] The above (a) to (e) will be described below.
(a) The ratio of the number of carbides at the ferrite grain
boundary relative to the number of carbides in the ferrite grain
exceeds 1:
[0089] The inventive steel sheet has a structure which is
substantially composed of ferrite and carbide, and in which the
ratio of the number of carbides at the ferrite grain boundary
relative to the number of carbides in the ferrite grain exceeds 1.
Carbides are, in addition to cementite (Fe.sub.3C) that is a
compound of iron and carbon, compounds obtained by replacing Fe in
cementite with an element such as Mn and Cr, and alloy carbides
(M.sub.23C.sub.6, M.sub.6Co, MC, etc., wherein M is Fe and another
additive metal element).
[0090] When a steel sheet is formed into a predetermined part
shape, a shear band is formed in the macrostructure of the steel
sheet, and slip deformation is generated and concentrated in the
vicinity of the shear band. The slip deformation involves
propagation of dislocations, and regions with high dislocation
density are formed in the vicinity of the shear band. As the strain
amount applied to the steel sheet increases, the slip deformation
is promoted and thereby the dislocation density increases. In cold
forging, strong processing exceeding an equivalent strain of 1 is
applied.
[0091] For this reason, in the conventional steel sheet, generation
of voids and/or cracks due to the increased dislocation density
could not be prevented, and it was difficult to improve cold
forgeability.
[0092] In order to solve the above challenging problems, it is
effective to suppress the formation of shear bands during forming.
From the viewpoint of a microstructure, shear band formation is a
phenomenon in which a slip generated in one crystal grain crosses
the crystal grain boundary and propagates continuously to an
adjacent crystal grain. Therefore, in order to suppress the
formation of a shear band, it is necessary to prevent the
propagation of slippage beyond crystal the grain boundary.
[0093] Carbides in the steel sheet are tenacious particles that
hinder slippage. Therefore, the presence of carbides at the ferrite
grain boundaries would make it possible, for the first time, to
suppress the formation of a shear band and thereby to improve cold
forgeability.
[0094] Based on the theory and principle, it is considered that
cold forgeability is strongly influenced by the coverage rate of
carbides at the ferrite grain boundaries. Therefore, it becomes
necessary to measure the coverage rate with high accuracy.
[0095] In order to measure the coverage rate of carbides at the
ferrite grain boundaries in a three-dimensional space, serial
sectioning SEM observation or repeated three-dimensional EBSP
observation become essential in which sample cutting by FIB and
observation are repeated in the scanning electron microscope.
However, these methods take a huge amount of measurement time and
the accumulation of technical know-how becomes indispensable. We
clarified this fact and concluded that common analytical methods
are not suitable.
[0096] Therefore, as a result of searching a simple and highly
accurate evaluation index, the present inventors have found that
cold forgeability can be evaluated by using, as an index, the ratio
of the number of carbides at the ferrite grain boundary relative to
the number of carbides in the ferrite grain, and that cold
forgeability can be remarkably improved when the ratio of the
number of carbides at the ferrite grain boundary relative to the
number of carbides in the ferrite grain is more than 1.
[0097] Any of buckling, folding and convolution of a steel sheet
that occurs during cold working is caused by the localization of
strain accompanying the formation of a shear band. Therefore, by
allowing the carbide to exist at the ferrite grain boundaries, the
formation of the shear band and the localization of strain can be
alleviated and the generation of buckling, folding and convolution
can be suppressed.
[0098] When the spheroidization percentage of carbides on the
crystal grain boundary is less than 80%, strains are concentrated
locally on the rod-shaped or plate-shaped carbides, and voids
and/or cracks are likely to occur. Therefore, the carbide
spheroidization ratio on the crystal grain boundary may preferably
be 80% or more, and more preferably 90% or more.
[0099] When the average particle diameter of the carbide in the
ferrite grain and the carbide at the ferrite grain boundaries is
less than 0.1 .mu.m, the hardness of the steel sheet remarkably
increases and the workability deteriorates. Therefore, the average
particle diameter of the carbide may preferably be 0.1 .mu.m or
more, and more preferably 0.17 .mu.m or more. On the other hand,
when the average particle diameter of the carbide exceeds 2.0
.mu.m, fissures occur with the coarse carbide serving as a starting
point during cold working, and thus the cold workability
deteriorates. Therefore, the average particle diameter of the
carbide may preferably be 2.0 .mu.m or less, and more preferably
1.95 .mu.m or less.
[0100] Subsequently, the method of observing and measuring the
structure will be described.
[0101] Observation of the carbide is carried out by a scanning
electron microscope. Prior to observation, samples for structure
observation are polished by wet polishing with emery paper and
polishing with diamond abrasive grains having an average particle
size of 1 .mu.m. After polishing the observation surface to a
mirror finish, the structure is etched with a 3% nitric
acid-alcohol solution.
[0102] Among the magnification for observation, within 3000 times,
a magnification capable of discriminating between ferrite and
carbide is selected. At the selected magnification, eight images
with a viewing field of 30 .mu.m.times.40 .mu.m are randomly
photographed at the 1/4 plate layer thickness.
[0103] With respect to the tissue image obtained, the area of each
carbide contained in the region is measured in detail by an image
analysis software represented by Mitsuya Shoji Co. Ltd. (Win ROOF).
A circle equivalent diameter (=2.times. (area/3.14)) is obtained
from the area of each carbide, and the average value is taken as
the carbide particle diameter.
[0104] Further, the spheroidization ratio of the carbide was
determined by approximating the carbide to an ellipse having an
equal area and equal moment of inertia, and then by calculating the
proportion of the carbides in which the ratio of the maximum length
to the maximum length in the perpendicular direction is less than
3.
[0105] In order to suppress the effect of measurement error due to
noise, carbides having an area of 0.01 .mu.m.sup.2 or more among
the carbides in grains and grain boundaries were counted and the
carbides having an area of 0.01 .mu.m.sup.2 or less were excluded
from evaluation.
[0106] The number of carbides present on the ferrite grain boundary
was counted, and from the total number of carbides the number of
carbides in the ferrite grain was determined by subtracting the
number of carbides on the ferrite grain boundary. Based on the
measured number, the ratio of the number of carbides on the grain
boundary relative to the number of carbides in the ferrite grain
was determined.
[0107] (b) The ferrite grain diameter is 5 .mu.m or more and 50
.mu.m or less:
[0108] In the structure after annealing the cold rolled steel
sheet, the cold workability can be improved by setting the ferrite
grain diameter to 5 .mu.m or more. When the ferrite grain size is
less than 5 .mu.m, the hardness increases and fissures and cracks
tend to generate easily during cold working. Therefore, the ferrite
grain size is set to 5 .mu.m or more, and preferably 7 .mu.m or
more.
[0109] On the other hand, when it exceeds 50 .mu.m, the number of
carbides on the crystal grain boundary that suppress slippage
propagation decreases and the cold workability deteriorates.
Therefore, that the ferrite grain size is set to 50 .mu.m or less,
and preferably 37 .mu.m or less.
[0110] The ferrite grain diameter is measured in the
above-described polishing method, wherein the observation surface
of the sample is polished to a mirror surface, followed by etching
with a 3% nitric acid-alcohol solution. The structure of the
observation surface is then examined with an optical microscope or
a scanning electron microscope, and a line segment method is then
applied to the image photographed to determine the ferrite grain
diameter.
[0111] (c) The in-plane anisotropy |.DELTA.r| of the r value
standardized according to JIS Z 2254 is 0.2 or less:
[0112] The in-plane anisotropy |.DELTA.r| of the plastic strain
ratio (r value) of the steel sheet is measured in a method in
accordance with JIS Z 2254. The r value (0.degree. direction:
r.sub.0, 45.degree. direction: r.sub.45, 90.degree. direction:
r.sub.90) measured by taking test strips from each direction of
0.degree. direction, 45.degree. direction and 90.degree. direction
with respect to the rolling direction was used to calculate the
following equation.
|.DELTA.r|=(r.sub.0-2r.sub.45+r.sub.90)/2
[0113] By setting the in-plane anisotropy |.DELTA.r| of the plastic
strain ratio (r value) of the steel sheet to 0.2 or less, the cold
workability can be improved. When |.DELTA.r| exceeds 0.2, the
thickness of parts and the height of the earing become uneven
during drawing. Therefore, the in-plane anisotropy |.DELTA.r| is
set to 0.2 or less.
[0114] (d) The Vickers hardness is 100 HV or more and 150 HV or
less:
[0115] By setting the Vickers hardness of the steel sheet to 100 HV
or more and 150 HV or less, the cold workability can be improved.
When the Vickers hardness is less than 100 HV, buckling can easily
occur during cold working. Therefore, the Vickers hardness is set
to 100 HV or more, and preferably 110 HV or more.
[0116] On the other hand, when the Vickers hardness exceeds 150 HV,
the ductility decreases and the internal breaking tends to occur
easily during cold forging. Therefore, the Vickers hardness is set
to 150 HV or less, and preferably 146 HV or less.
[0117] (e) The ratio of X-ray diffraction intensity of the {311}
<011> orientation at the 1/2-thickness portion of the steel
sheet relative to the X-ray diffraction intensity obtained when a
sample with a random orientation distribution of crystal grains in
the steel sheet is subjected to X-ray diffraction is 3.0 or
less:
[0118] In cold forging, in addition to controlling the figuration
of carbides, the draw formability during cold forging must be
secured. In order to improve the draw formability during cold
forging, plastic anisotropy such as in-plane anisotropy |.DELTA.r|
must be improved. For that purpose, the aggregate structure of a
hot-rolled steel sheet must be controlled. For evaluation of the
aggregate structure, analysis by X-ray diffraction on a plane
parallel to the plate surface at the 1/2 thickness portion of the
hot-rolled steel plate is used.
[0119] One surface of a hot-rolled steel plate is ground to a 1/2
plate thickness surface in parallel to the surface to expose a 1/2
plate thickness surface, followed by the analysis of the 1/2 plate
thickness surface by X-ray diffraction. As the X-ray diffraction,
X-ray diffraction by Mo bulb may be used. Diffraction intensities
of diffraction orientations {110}, {220}, {211} and {310} by
reflection are obtained, and based thereon, the orientation
distribution function (ODF) is created.
[0120] The X-ray diffraction intensity ratio is determined by using
the diffraction intensity data of the 1/2 plate thickness surface
obtained from the ODF and the diffraction intensity data of random
orientation of the hot-rolled steel sheet. Specifically, as a
standard sample in which the metallic structure has no accumulation
in a specific direction, a sample obtained by sintering powder iron
of a hot-rolled steel sheet to be measured or the powder before
sintering is used to determine the diffraction intensity under the
same conditions as when the diffraction intensity data of the 1/2
plate thickness surface was obtained. The part to be collected as
the standard sample is not particularly limited and may be any part
of the hot-rolled steel sheet. The X-ray diffraction intensity
ratio in a specific orientation is a numerical value obtained by
dividing the diffraction intensity in the specific direction of the
1/2 plate thickness surface obtained from the ODF by the
diffraction intensity of the standard sample.
[0121] When the X-ray diffraction intensity ratio of the {311}
<011> orientation obtained by the above-described ODF
analysis is set to I1, it is necessary that this I1 is 3.0 or less,
and preferably 2.5 or less for the random aggregate structure
during hot rolling. When a random aggregate structure having I1 of
3.0 or less can be obtained, the plasticity anisotropy is reduced
and the cold formability is improved.
[0122] Next, the inventive manufacturing method will be
described.
[0123] The manufacturing method according to the present invention
is characterized in that the hot rolling and the annealing are
consistently managed to control the structure. After continuously
casting a steel strip having a predetermined ingredient
composition, the steel strip is subjected to hot rolling by heating
to complete finish hot rolling at a temperature range of
800.degree. C. or higher to 900.degree. C. or lower, coiled at
400.degree. C. or higher and 550.degree. C. or lower to obtain a
hot-rolled steel sheet. The hot-rolled steel sheet is, after
pickling, subjected to a two-step type annealing in which the
hot-rolled steel sheet is maintained in two temperature ranges,
whereupon
[0124] (i) the hot-rolled steel sheet is subjected to a first step
annealing performed by retaining said hot-rolled steel at a
temperature range of 650.degree. C. or higher and 720.degree. C. or
lower for 3 hours or longer and 60 hours or shorter, and then
subjected to a second step annealing performed by retaining the
hot-rolled steel at a temperature range of 725.degree. C. or higher
and 790.degree. C. or lower for 3 hours or longer and 50 hours or
shorter, and thereafter
[0125] (ii) the hot-rolled steel sheet is cooled to 650.degree. C.
or lower at a cooling rate of 1.degree. C./hour or more and
30.degree. C./hour or less,
[0126] and thus a steel sheet excellent in cold workability during
forming can be produced.
[0127] By the hot rolling and annealing mentioned above, a
structure composed of fine pearlite and bainite can be formed as
the structure of the steel sheet.
[0128] The processing conditions will be described below.
Heating temperature of a steel strip: 1000.degree. C. or higher and
1250.degree. C. or lower
[0129] The heating temperature of the steel strip subjected to hot
rolling may preferably be 1000.degree. C. or higher and
1250.degree. C. or lower, and the heating time may preferably be
0.5 hour or longer and 3 hours or shorter.
[0130] When the heating temperature is lower than 1000.degree. C.
or the heating time is shorter than 0.5 hour, the microsegregation
and/or macrosegregation formed by casting are not eliminated, and
regions in which Si, Mn, etc., are locally concentrated inside the
steel material may remain, and thus the impact resistance property
of the steel material is lowered. Therefore, the heating
temperature may preferably be 1000.degree. C. or higher, and
preferably 0.5 hour or longer.
[0131] On the other hand, when the heating temperature exceeds
1250.degree. C. or the heating time exceeds 3 hours,
decarburization from the surface layer of the steel strip becomes
conspicuous, and austenite grains in the surface layer grow
abnormally during heating before carburizing and quenching, and the
impact resistance property of the steel strip is deteriorated. Thus
the heating temperature may preferably be 1250.degree. C. or lower,
and the heating time may preferably be 3 hours or shorter.
Finish hot rolling temperature: 800.degree. C. or higher and
900.degree. C. or lower
[0132] Finish hot rolling is completed at 800.degree. C. or higher
and 900.degree. C. or lower. When the finish hot rolling
temperature is lower than 800.degree. C., the deformation
resistance of the steel strip increases, the rolling load increases
markedly, the wear amount of the roll increases, and the
productivity decreases. Therefore, the finish hot rolling
temperature is set to 800.degree. C. or higher, and preferably
820.degree. C. or higher.
[0133] On the other hand, when the finish hot rolling temperature
exceeds 900.degree. C., thick scales are generated during plate
passing on the ROT (Run Out Table), scratches are generated on the
surface of the steel sheet due to the scale, and cracks are
generated starting from scratches when an impact load is applied
after cold forging and carburizing and annealing, leading to
reduced impact resistance property of the steel sheet. Therefore,
the finish hot rolling temperature is set to 900.degree. C. or
lower, and preferably 880.degree. C. or lower.
Cooling rate on ROT: 10.degree. C./sec or more and 100.degree.
C./sec or less
[0134] The cooling rate at the time of cooling the hot-rolled steel
sheet on the ROT after finish hot rolling may preferably be
10.degree. C./sec or more and 100.degree. C./sec or less. When the
cooling rate is less than 10.degree. C./sec, thick scales are
generated during cooling and the occurrence of scratches on the
surface of the steel sheet due to the scales cannot be suppressed.
Therefore, the cooling rate is set to 10.degree. C./sec or more,
and more preferably 20.degree. C./sec or more.
[0135] On the other hand, when the cooling rate exceeds 100.degree.
C./sec, the steel sheet is cooled at a cooling rate exceeding
100.degree. C./sec from the surface layer to the inside of the
steel sheet, the outermost layer part of the steel sheet is
excessively cooled, and a low-temperature transformed structure
such as bainite or martensite is formed.
[0136] At the time of discharging the hot-rolled coil cooled from
100.degree. C. to room temperature after coiling, microcracks are
generated in the low-temperature transformed structure. It is
difficult to remove the microcracks in the subsequent pickling step
and cold rolling step, and fissures progress from the microcracks
as a starting point during cold working, leading to reduced cold
workability. Therefore, the cooling rate may preferably be
100.degree. C./sec or less.
[0137] Note that the above cooling rate refers to the cooling
capacity from the cooling facility at each water injection zone
from the point at which the hot-rolled steel sheet after the finish
hot rolling is cooled at the water injection zone after passing
through the water-free zone to a point at which it is cooled to the
coiling target temperature on the ROT, and does not refer to the
average cooling rate from the water injection starting point to the
temperature at which it is coiled by the coiling device.
Coiling temperature: 400.degree. C. or higher and 550.degree. C. or
lower
[0138] The coiling temperature is set to 400.degree. C. or higher
and 550.degree. C. or lower. When the coiling temperature is lower
than 400.degree. C., the austenite which was not transformed before
coiling is transformed into hard martensite, cracks are generated
in the surface layer of the steel sheet during discharge of the
hot-rolled coil, leading to reduced workability. Therefore, the
coiling temperature is set to 400.degree. C. or higher, and
preferably 430.degree. C. or higher.
[0139] On the other hand, when the coiling temperature exceeds
550.degree. C., pearlite having a large lamellar spacing is
generated and thick needle-shaped carbides having high thermal
stability are formed, and even after the two-step type annealing,
needle-shaped carbides remain. Since fissures are generated during
cold working with these needle-shaped carbides as a starting point,
the coiling temperature is set to 550.degree. C. or lower, and
preferably 520.degree. C. or lower.
[0140] The hot-rolled coil manufactured under the above conditions
is annealed, after pickling, in a two-step type annealing which
retains the coil in two temperature ranges. The first-step
annealing and the second-step annealing may be either box annealing
or continuous annealing. By controlling the stability of carbides
by the two-step type annealing, the formation of carbides on the
ferrite grain boundary and the spheroidization ratio of carbides on
the ferrite grain boundary can be enhanced.
[0141] The two-step type annealing will be described below.
[0142] The first step annealing is carried out in a temperature
range of the A.sub.ci point or lower to coarsen carbides and enrich
alloy elements to increase the thermal stability of carbides.
Thereafter, the temperature is raised to a range from A.sub.C1
point or higher to A.sub.3 point or lower to generate austenite in
the structure.
[0143] Thereafter, by gradual cooling, the austenite is transformed
into ferrite and the carbon concentration in the austenite is
increased. By proceeding slow cooling, carbon atoms are adsorbed to
the carbides remaining in the austenite, and thus the carbide and
austenite come to cover the grain boundary of the ferrite. Finally
it becomes possible to form a structure in which many spheroidized
carbides are present in the grain boundary of the ferrite.
[0144] When the residual carbides are small in quantity while
maintaining the temperature range of A.sub.C1 point or higher to
A.sub.3 point or lower, pearlite, rod-shaped carbides and
plate-like carbides are produced during cooling. When these
pearlite, rod-shaped carbides and plate-like carbides are produced,
the workability of the steel sheet is remarkably deteriorated.
Therefore, increasing the number of residual carbides in the
temperature range from A.sub.C1 point or higher to A.sub.3 point or
lower is an important factor to enhance the workability of the
steel sheet.
[0145] By using a steel sheet structure obtained under the above
hot rolling condition, the thermal stability of carbides at a
temperature of A.sub.C1 point or lower can be secured. Therefore,
an increase in the number of residual carbides in the temperature
range from A.sub.C1 point or higher to A.sub.3 point or lower can
be targeted.
[0146] Hereinafter, an annealing condition for the two-step type
annealing will be described.
First step annealing
[0147] Temperature range: 650.degree. C. or higher and 720.degree.
C. or lower
[0148] Retention time: 3 hours or longer and 60 hours or
shorter
[0149] In the first step annealing, the annealing temperature is
set to 650.degree. C. or higher and 720.degree. C. or lower. When
the annealing temperature of the first step is lower than
650.degree. C., the stability of the carbide becomes insufficient
and it becomes difficult to allow the carbide to remain in the
austenite in the second step annealing. Therefore, the temperature
of the first step annealing is set to 650.degree. C. or higher, and
preferably 670.degree. C. or higher.
[0150] On the other hand, when the temperature of the first step
annealing exceeds 720.degree. C., austenite is generated before
enhancing the stability of the carbide, which makes it difficult to
control the required change in the structure. Therefore, the first
step annealing temperature is set to 720.degree. C. or lower, and
preferably 700.degree. C. or lower.
[0151] The retention time at the first step is 3 hours or longer
and 60 hours or shorter. When the retention time is lower than 3
hours, the stability of the carbide is insufficient and it becomes
difficult to allow the carbide to remain at the second step
annealing. Therefore, the retention time of the first step is set
to 3 hours or longer. On the other hand, when the retention time of
the first step exceeds 60 hours, improvement of the stability of
the carbide cannot be expected and furthermore the productivity is
lowered. Therefore, the retention time of the first step is set to
60 hours or shorter, and preferably 55 hours or shorter.
[0152] The annealing atmosphere is not limited to a specific
atmosphere. For example, it may be either a nitrogen atmosphere
having a nitrogen content of 95% or more, a hydrogen atmosphere
having a hydrogen content of 95% or more, or an atmospheric
atmosphere.
Second step annealing
[0153] Temperature range: 725.degree. C. or higher and 790.degree.
C. or lower
[0154] Retention time: 3 hours or longer and 50 hours or
shorter
[0155] In the second step annealing, the annealing temperature is
set to 725.degree. C. or higher and 790.degree. C. or lower. When
the second-step annealing temperature is lower than 725.degree. C.,
the amount of austenite produced is small and the number ratio of
carbides on the ferrite grain boundary is lowered. Therefore, the
second-step annealing temperature is set to 725.degree. C. or
higher, and preferably 745.degree. C. or higher.
[0156] On the other hand, when the second-step annealing
temperature exceeds 790.degree. C., it becomes difficult to allow
the carbide to remain in the austenite and to control the required
structure change. Therefore, the second-step annealing temperature
is set to 790.degree. C. or lower, and preferably 770.degree. C. or
lower.
[0157] The retention time of the second step is set to 3 hours or
longer and 50 hours or shorter. When the retention time of the
second step is less than 3 hours, the amount of austenite produced
is small, dissolution of the carbide in the ferrite grains is
insufficient, and it becomes difficult to increase the number ratio
of carbides on the ferrite grain boundary. Therefore, the retention
time of the second step is set to 3 hours or longer, and preferably
5 hours or longer.
[0158] On the other hand, when the retention time of the second
step exceeds 50 hours, it becomes difficult to allow the carbide to
remain in the austenite. Therefore, the retention time of the
second step is set to 50 hours or shorter, and preferably is 46
hours or shorter.
[0159] The annealing atmosphere is not limited to a specific
atmosphere. For example, it may be either a nitrogen atmosphere
having a nitrogen content of 95% or more, a hydrogen atmosphere
having a hydrogen content of 95% or more, or an atmospheric
atmosphere.
[0160] After completion of the two-step type annealing, the
hot-rolled steel sheet is cooled, whereupon it is cooled to
650.degree. C. at a cooling rate of 1.degree. C./hour or more to
30.degree. C./hour or less.
[0161] Cooling rate to a temperature of 650.degree. C. or lower:
1.degree. C./hour or more and 30.degree. C./hour or less
[0162] Since the temperature range for controlling the structure
change by slow cooling is sufficient up to 650.degree. C., it is
only necessary to control the cooling rate in the temperature range
up to 650.degree. C. After reaching a temperature of 650.degree. C.
or lower, it may be cooled to room temperature within the above
range without controlling the cooling rate.
[0163] It may be preferable that the cooling rate is slow in order
to gradually cool the austenite produced in the second step
annealing to transform into ferrite and allow carbon to be adsorbed
to the carbides remaining in the austenite. However, when the
cooling rate is less than 1.degree. C./hour, the time required for
cooling increases and the productivity decreases. Therefore, the
cooling rate is 1.degree. C./hour or more, and preferably 5.degree.
C./hour.
[0164] On the other hand, when the cooling rate exceeds 30.degree.
C./hour, austenite transforms to pearlite, the hardness of the
steel sheet increases, the cold forgeability deteriorates, and the
impact resistance property of the steel sheet after carburizing
quenching and tempering decreases. Therefore, the cooling rate is
set to 30.degree. C./hour or less, and preferably 26.degree.
C./hour or less.
[0165] Further, according to the inventive production method, a
steel sheet with excellent cold workability during forming can be
produced in which the ingredient composition is, in terms of % by
mass, comprising: C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to
1.00% , P: 0.0001 to 0.020%, S: 0.0001 to 0.010%, and Al: 0.001 to
0.10%, the balance being Fe and unavoidable impurities, the metal
structure is substantially composed of ferrite and spheroidized
carbides, and (a) the ratio of the number of carbides at the
ferrite grain boundary to the number of carbides in the ferrite
grain exceeds 1, (b) the ferrite grain size is 5.mu.m or more and
50 .mu.m or less, (c) the in-plane anisotropy |.DELTA.r| of the r
value standardized according to JIS Z 2254 is 0.2 or less, (d) the
Vickers hardness is 100 HV or more and 150 HV or less, the
cross-sectional shrinkage percentage is 40% or more, and the ratio
of X-ray diffraction intensity of the {131} <011> orientation
at the 1/2-thickness portion of the steel sheet relative to the
X-ray diffraction intensity obtained when a sample having the
random orientation distribution of crystal grains in the steel
sheet is subjected to X-ray diffraction is 3.0 or less.
[0166] The cross-sectional shrinkage percentage is defined by the
following formula (1). A large value of this ratio means that the
local deformability is high, and as the value of the formula (1)
increases, the workability of the steel sheet increases.
Sectional shrinkage percentage (%)=100-(cross-sectional area at
tensile fracture/initial cross-sectional area).times.100 Equation
(1)
[0167] As described above, the present invention is characterized
in that by rolling control and heat treatment after rolling, a
structure in which carbides (that is, cementite) are uniformly
dispersed is formed, so that the crystal anisotropy can be
eliminated. Therefore, in the present invention, the random
intensity ratio of the {311} <011> orientation at the 1/2
plate thickness portion of the steel sheet can be made 3.0 or
less.
EXAMPLES
[0168] Next, examples will be described, but the level of examples
is an example of conditions adopted for confirming the feasibility
and effectiveness of the present invention, and the present
invention is not limited to this one condition example. The present
invention can adopt various conditions as long as the object of the
present invention is achieved without departing from the gist of
the present invention.
Example 1
[0169] In order to investigate the effect of hot rolling
conditions, a continuous cast strip (steel ingot) having the
ingredient composition shown in Table 1 was subjected to hot
rolling under the conditions shown in Table 2 to produce a
hot-rolled coil having a thickness of 3.0 mm. Incidentally, the
steel type described as "Developed steel" in the column of
"Remarks" in Table 1 has a composition included in the composition
range of the steel sheet according to the present invention. Also,
the steel type described as "Comparative steel" in the column of
"Remarks" in Table 1 has a composition outside the composition
range of the steel sheet according to the present invention. In
addition, the ingredients that do not satisfy the composition
conditions of the steel sheet according to the present invention
are underlined.
[0170] A sample for characterization was prepared as follows: a
hot-rolled coil, after pickling, was placed in a box-type annealing
furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen,
the coil was heated from room temperature to 705.degree. C. and was
retained for 36 hours to make the temperature distribution uniform
in the hot-rolled coil. The coil was then heated to 760.degree. C.
and retained at 760.degree. C. for 10 hours, and was then cooled to
650.degree. C. at a cooling rate of 10.degree. C./hour, then
furnace-cooled to room temperature to prepare the sample for
characterization. The structure of the sample was measured by the
method described above.
TABLE-US-00001 TABLE 1 C Si Mn P S Al N O Cr Mo Nb V Cu W A 0.21
0.13 0.53 0.0048 0.0084 0.618 0.0053 0.098 0.084 B 0.24 0.28 0.91
0.0039 0.0018 C 0.34 0.19 0.51 0.0133 0.0015 0.049 D 0.18 0.18 0.50
0.0151 0.0092 0.0081 0.461 E 0.32 0.03 0.73 0.0149 0.0094 0.018
0.0018 0.087 0.080 F 0.22 0.06 0.98 0.0144 0.0088 0.082 0.073 0.039
G 0.17 0.13 0.94 0.0199 0.0026 0.038 0.277 H 0.37 0.11 0.38 0.0144
0.0087 I 0.33 0.10 0.96 0.0010 0.0019 1.180 0.480 J 0.29 0.16 0.81
0.0132 0.0031 K 0.11 0.02 0.41 0.0117 0.0360 L 0.25 0.11 0.63
0.0167 0.0058 0.0007 0.090 0.020 M 0.19 1.21 0.91 0.0027 0.0084 N
0.31 0.23 1.27 0.0010 0.0019 O 0.51 0.14 0.65 0.0059 0.0013 P 0.20
0.13 0.92 0.0153 0.0091 Q 0.11 0.07 0.64 0.0060 0.0051 0.0013
0.0131 R 0.27 0.19 0.70 0.0185 0.0033 0.012 S 0.27 0.13 0.80 0.0044
0.0002 0.071 0.060 T 0.25 0.14 0.38 0.0142 0.0097 U 0.12 0.25 0.86
0.0154 0.0045 W 0.40 0.08 0.78 0.0184 0.0087 0.026 0.025 0.088
0.055 X 0.13 0.20 0.96 0.0061 0.0034 Y 0.21 0.21 0.41 0.0139 0.0038
0.092 0.093 Z 0.11 0.05 0.59 0.0195 0.0010 AA 0.13 0.09 0.76 0.0127
0.0079 AB 0.19 0.22 0.89 0.0005 0.0021 AC 0.12 0.02 0.63 0.0168
0.0098 AD 0.30 0.28 0.71 0.0169 0.0031 0.018 0.035 Ta Ni Sn Sb As
Mg Ca Y Zr La Ce Remarks A 0.004 0.015 0.0149 Comparative steel B
Developed steel C 0.083 0.018 0.0459 Developed steel D 0.017 0.018
Developed steel E 0.020 0.034 0.021 0.045 Developed steel F
Developed steel G 0.014 Developed steel H Developed steel I 0.037
0.046 0.048 Comparative steel J Developed steel K Comparative steel
L 0.012 0.0423 0.035 0.008 Developed steel M 0.037 0.046 0.048
Comparative steel N Comparative steel O Comparative steel P
Developed steel Q 0.033 0.035 Developed steel R 0.080 0.037
Developed steel S 0.016 0.0280 0.036 Developed steel T Developed
steel U Developed steel W 0.038 0.015 Developed steel X Developed
steel Y 0.039 0.030 Developed steel Z Developed steel AA Developed
steel AB Developed steel AC Developed steel AD 0.041 Developed
steel
TABLE-US-00002 TABLE 2 Hot rolling Condition Grain Cross-
{311}<011> Finish hot Ferrite boundary sectional X-ray
rolling Coiling Carbide grain Vickers carbide shrinkage diffraction
In-plane temp. temp. diameter diameter hardness No./grain percent
intensity anisotropy [.degree. C.] [.degree. C.] [.mu.m] [.mu.m]
[HV] carbide No. [%] ratio (I1) |.DELTA.r| Remarks A-1 891 505 0.91
18.5 125 3.7 51.0 5.5 0.44 Comparative steel B-1 832 500 1.17 20.7
124 6.7 49.7 1.6 0.05 Inventive steel C-1 850 489 1.08 20.7 120 4.6
51.3 2.1 0.10 Inventive steel D-1 742 532 0.99 21.5 111 6.3 54.0
5.6 0.48 Comparative steel E-1 811 617 0.95 16.8 116 5.4 29.7 1.5
0.06 Comparative steel F-1 937 512 1.18 21.1 112 7.2 52.8 2.5 0.18
Comparative steel G-1 871 409 1.27 23.9 110 8.7 56.3 1.5 0.04
Inventive steel H-1 837 522 0.99 20.7 117 3.7 55.7 2.0 0.10
Inventive steel I-1 802 484 1.03 16.9 121 7.1 49.5 5.0 0.39
Comparative steel J-1 841 407 1.28 22.5 116 6.9 43.9 1.9 0.09
Inventive steel K-1 871 475 1.11 29.7 104 4.4 31.4 1.4 0.06
Comparative steel L-1 898 443 1.20 22.7 110 6.9 54.7 2.3 0.12
Inventive steel M-1 873 467 1.11 18.1 172 5.4 34.9 4.6 0.37
Comparative steel N-1 850 414 1.37 22.8 152 9.3 38.5 2.4 0.14
Comparative steel O-1 834 413 1.22 20.0 127 5.7 38.4 2.0 0.12
Comparative steel P-1 856 485 1.22 22.8 111 7.2 57.1 2.2 0.11
Inventive steel Q-1 873 428 1.25 29.3 116 6.2 54.1 2.0 0.12
Inventive steel R-1 845 348 1.32 24.0 114 6.5 51.1 2.2 0.14
Comparative steel S-1 837 454 1.23 21.9 114 7.1 56.6 1.8 0.08
Inventive steel T-1 830 671 0.75 17.4 116 2.8 38.7 1.4 0.05
Comparative steel U-1 881 419 1.29 27.6 112 7.3 59.0 2.7 0.19
Inventive steel W-1 853 535 1.09 18.5 121 6.1 55.1 1.4 0.05
Inventive steel X-1 885 477 1.24 25.8 112 7.5 58.0 2.9 0.19
Inventive steel Y-1 874 436 1.14 25.7 110 5.8 44.8 2.0 0.12
Inventive steel Z-1 899 485 1.16 27.7 108 5.5 57.5 1.6 0.06
Inventive steel AA-1 895 423 1.28 27.6 102 6.9 42.6 2.4 0.14
Inventive steel AB-1 855 434 1.27 23.9 115 7.3 53.3 2.5 0.16
Inventive steel AC-1 895 381 1.30 30.1 109 6.4 54.4 1.9 0.11
Comparative steel AD-1 854 463 1.17 20.7 125 5.9 52.3 1.4 0.06
Inventive steel
[0171] The cold workability was evaluated using the notched tensile
test and the in-plane anisotropy of the r value. In the notched
tensile test, a notched tensile test strip was taken from an
as-annealed material with a thickness of 3 mm, and a tensile test
was performed in the rolling direction to determine the
cross-sectional shrinkage percentage, and the local deformability
was evaluated. When the cross-sectional shrinkage percentage is 40%
or more, it was rated as superior.
[0172] Further, the in-plane anisotropy of the r value was rated as
superior when the in-plane anisotropy |.DELTA.r| of the r value
standardized according to JIS Z 2254 of an as-annealed material
with a thickness of 3 mm was 0.2 or less.
[0173] In order to determine the X-ray diffraction intensity ratio
(I1) of {311} <011>, X-ray diffraction with an Mo tube was
performed from the center of the plate thickness of each sample
followed by an ODF analysis. Based on the results obtained by the
ODF analysis, the I1 was determined.
[0174] Table 2 shows, for each of the samples prepared, the results
of the carbide diameter, the ferrite grain diameter, the Vickers
hardness, the ratio of the number of carbides at the ferrite grain
boundary relative to the number of carbides in the ferrite grain,
the cross-sectional shrinkage percentage, the X-ray diffraction
intensity ratio of {311} <011> and in-plane anisotropy. Among
the samples in Table 2, those indicated as "Inventive steel" in the
Remarks column satisfy the requirements of the steel sheet
according to the present invention, and those indicated as
"Comparative steel" in the Remarks column do not satisfy the
requirements of the steel sheet according to the present invention.
In Table 2, the measurement results that do not satisfy the
requirements of the steel sheet according to the present invention
and the manufacturing conditions that do not satisfy the
requirements of the steel sheet manufacturing method according to
the present invention are underlined.
[0175] As shown in Table 2, in any of the inventive steels B-1,
C-1, G-1, H-1, J-1, L-1, P-1, Q-1, S-1, U-1, W-1, X-1, Y-1, Z-1,
AA-1, AB-1 and AD-1, the ratio of the number of carbides at the
ferrite grain boundary relative to the number of carbides in the
ferrite grain exceeds 1, and the Vickers hardness is 150 HV or
less. In addition, in any of the inventive steels, the
cross-sectional shrinkage percentage exceeds 40% and the in-plane
anisotropy |.DELTA.r| of the r value is 0.2 or less. Thus, they
have excellent cold workability. Furthermore, since it was
confirmed that scale scratches were not generated on the steel
sheet surface in any of the inventive steels, these steels can be
suitably used for cold working.
[0176] On the other hand, in the Comparative steel A-1, since the
Al content is high and the A3 point decreased, recrystallization
during finish hot rolling was inhibited and |.DELTA.r|
deteriorated. Thus, the cold workability is low. In the Comparative
steel I-1, the contents of Mo and Cr are high, recrystallization
during finish hot rolling was inhibited, and |.DELTA.r|
deteriorated. In the comparative steels K-1 and N-1, the content of
S or Mn is high, coarse MnS was formed in the steel, and the cold
workability is low. In the Comparative steel M-1, the content of Si
was high and hardness increased, and thus cold workability is low.
Also, in the Comparative steel M-1, since the A3 point rose,
recrystallization during finish hot rolling was hindered and
|.DELTA.r| deteriorated.
[0177] In the Comparative steel O-1, C is high, the volume fraction
of carbides increased, a large amount of cracks as the starting
point of fractures were generated, and the cross-sectional
shrinkage percentage was low. Thus, the cold workability is low. In
the Comparative steel D-1, the finish temperature of hot rolling
was low and the productivity decreased. In the Comparative steel
F-1, the finish temperature of hot rolling was high, and scale
scratches were generated on the surface of the steel sheet.
[0178] In the Comparative steels R-1 and AC-1, the coiling
temperature of hot rolling was low, the low-temperature
transformation structure such as bainite and martensite increased
resulting in brittled steel, and breaks frequently occurred when
the hot-rolled coil was discharged resulting in a decrease in
productivity. In the Comparative steels E-1 and T-1, the coiling
temperature of hot rolling was high, thick pearlite with lamellar
spacing and needle-shaped coarse carbides with high thermal
stability were produced in the hot rolled structure. Since these
carbides remained in the steel sheet even after the two-step type
annealing, the cross-sectional shrinkage percentage was low and
thus the cold workability is low.
[0179] Subsequently, in order to investigate the effect of
annealing conditions, steel strips (slabs) having the ingredient
composition shown in Table 1 were heated at 1240.degree. C. for 1.8
hours and then subjected to hot rolling. After completing finish
hot rolling at 890.degree. C., they were cooled to 520.degree. C.
at a cooling rate of 45.degree. C./sec on ROT and coiled at
510.degree. C. to produce a hot-rolled coil with a thickness of 3.0
mm. And under the conditions shown in Table 3, a hot-rolled
sheet-annealed sample with a thickness of 3.0 mm was prepared.
[0180] For each of the samples prepared, the carbide diameter, the
ferrite grain diameter, the Vickers hardness, the ratio of the
number of carbides at the ferrite grain boundary relative to the
number of carbides in the ferrite grain, the cross-sectional
shrinkage percentage, the X-ray diffraction intensity ratio of
{311} <011> and the in-plane anisotropy were determined in
the same manner as the inventive steels and the comparative steels
in Table 2. The results are shown in Table 3.
TABLE-US-00003 TABLE 3 1st step annealing 2nd step annealing
Ferrite Retention Retention Retention Retention Cooling Carbide
grain temp. time temp. time rate diameter diameter [.degree. C.]
[hr] [.degree. C.] [hr] [.degree. C./sec] [.mu.m] [.mu.m] A-2 669
27 761 41 7 0.96 27.1 B-2 695 47 753 12 30 0.80 15.2 C-2 654 19 771
32 20 0.90 23.3 D-2 705 51 760 23 13 0.98 25.4 E-2 693 50 729 45 34
0.60 13.6 F-2 698 22 743 36 14 0.91 16.8 G-2 695 55 759 19 7 1.26
24.2 H-2 698 30 742 1 6 0.98 13.5 I-2 658 57 754 14 16 0.75 12.8
J-2 694 42 811 29 7 1.59 38.1 K-2 675 21 746 42 22 0.71 26.5 L-2
694 14 781 12 30 0.84 22.0 M-2 709 9 779 8 29 0.81 18.2 N-2 670 15
738 14 25 0.76 10.9 O-2 676 52 732 46 7 0.82 14.6 P-2 671 21 769 42
9 1.23 29.5 Q-2 701 2 756 34 25 0.79 24.9 R-2 663 52 729 15 17 0.61
10.3 S-2 681 56 779 4 7 1.41 27.3 T-2 741 47 774 9 25 1.29 32.6 U-2
672 12 741 47 25 0.71 17.7 W-2 700 45 730 8 28 0.72 10.3 X-2 709 46
743 45 8 1.15 22.2 Y-2 676 51 777 54 17 0.94 37.7 Z-2 668 30 706 23
16 0.57 8.5 AA-2 662 14 776 26 19 2.23 54.0 AB-2 678 68 785 12 14
1.16 29.4 AC-2 637 35 745 48 6 1.19 26.2 AD-2 712 50 750 29 15 0.96
19.3 Grain {311}<011> boundary Cross- X-ray Vickers carbide
sectional diffraction In-plane hardness No./grain shrinkage
intensity anisotropy [HV] carbide No. ratio [%] ratio (I1)
|.DELTA.r| Remarks A-2 116 4.8 56.3 4.6 0.38 Comp. steel B-2 136
3.1 48.7 1.9 0.10 Developed steel C-2 112 2.8 58.7 1.8 0.11
Developed steel D-2 108 3.7 59.4 1.5 0.06 Developed steel E-2 161
1.7 38.6 1.2 0.04 Comp. steel F-2 120 4.9 54.8 2.5 0.17 Developed
steel G-2 118 11.4 51.0 2.2 0.15 Developed steel H-2 133 0.4 50.9
1.5 0.07 Comp. steel I-2 133 8.2 53.6 3.8 0.28 Comp. steel J-2 151
2.9 32.6 1.9 0.10 Comp. steel K-2 104 2.7 20.7 1.0 0.03 Comp. steel
L-2 109 4.3 55.9 2.4 0.17 Developed steel M-2 183 3.9 38.6 2.8 0.18
Comp. steel N-2 153 11.2 39.6 1.6 0.06 Comp. steel O-2 148 4.0 37.4
1.2 0.05 Comp. steel P-2 108 4.6 59.0 1.4 0.05 Developed steel Q-2
135 3.2 37.6 1.6 0.05 Comp. steel R-2 138 4.5 48.0 2.0 0.12
Developed steel S-2 126 7.6 49.9 1.2 0.04 Developed steel T-2 138
1.5 34.7 1.1 0.01 Comp. steel U-2 121 3.7 42.1 2.6 0.16 Developed
steel W-2 141 2.1 46.5 2.0 0.11 Developed steel X-2 118 4.2 49.6
1.8 0.08 Developed steel Y-2 141 2.8 29.2 1.8 0.07 Comp. steel Z-2
124 0.5 51.0 1.1 0.04 Comp. steel AA-2 105 9.4 60.2 1.4 0.05
Developed steel AB-2 116 4.8 45.7 1.4 0.06 Comp. steel AC-2 133 6.8
31.2 1.9 0.10 Comp. steel AD-2 127 2.3 50.9 1.7 0.06 Developed
steel
[0181] As shown in Table 3, in any of the inventive steels B-2,
C-2, D-2, F-2, G-2, L-2, P-2, R-2, S-2, U-2, W-2, X-2, AA-2 and
AD-2, the ratio of the number of carbides at the ferrite grain
boundary to the number of carbides in the ferrite grain exceeds 1,
and the Vickers hardness is 150 HV or less. In addition, in any of
the inventive steels, the cross-sectional shrinkage percentage
exceeds 40% and the in-plane anisotropy |.DELTA.r| of the r value
is 0.2 or less. Thus, they have excellent cold workability.
[0182] On the other hand, in the Comparative steel A-2, since the
Al content is high and the A3 point decreased, recrystallization
during finish hot rolling was inhibited and |.DELTA.r|
deteriorated. Thus, the cold workability is low. In the Comparative
steel 1-2, the contents of Mo and Cr are high, recrystallization
during finish hot rolling was inhibited, and |.DELTA.r|
deteriorated. In the comparative steels K-2 and N-2, the content of
S or Mn is high, coarse MnS was formed in the steel. Thus, the cold
workability deteriorated. In the Comparative steel M-2, the content
of Si was high and hardness increased. Thus, the cold workability
is low. Also, in the Comparative steel M-2, since the A3 point
decreased, recrystallization during finish hot rolling was hindered
and |.DELTA.r| deteriorated.
[0183] In the Comparative steel O-2, C is high, the volume fraction
of carbides increased, a large amount of cracks as the starting
point of fracture were generated, and the cross-sectional shrinkage
percentage was low. Thus, the cold workability is low.
[0184] In the Comparative steel AC-2, since the annealing
temperature in the first-step annealing during the two-step type
box annealing is low, the treatment of carbide coarsening at the
Ac1 temperature or lower is insufficient, and the thermal stability
of carbides is insufficient, thus the carbides remaining at the
second step of annealing decreases, the pearlite transformation
cannot be suppressed in the structure after the slow cooling, and
the cross-sectional shrinkage percentage is low. Thus, the cold
workability is low.
[0185] In the Comparative steel T-2, since the annealing
temperature in the first step annealing during the two-step type
box annealing is high, austenite is generated during annealing and
the stability of carbide cannot be increased, so that carbides
remaining during the second step annealing decrease, and pearlite
transformation cannot be suppressed in the structure after the slow
cooling, and the cross-sectional shrinkage percentage is low. Thus,
the cold forgeability is low.
[0186] In the Comparative steel Q-2, since the retention time in
the first step annealing during annealing of the two-step type is
short, the treatment of the carbide coarsening at the Ac1
temperature or lower is insufficient, and the thermal stability of
the carbide is insufficient, and thus the carbide remaining at the
second step of annealing decreases and the pearlite transformation
cannot be suppressed in the structure after the slow cooling, and
the cross-sectional shrinkage percentage is low. And thus the cold
workability is low. In the comparative steel AB-2, the retention
time during the first stage box annealing of the two-step type is
long and the productivity is low.
[0187] In the Comparative steel Z-2, since the annealing
temperature during the second-step annealing during the two-step
type box annealing is low, and the amount of austenite produced is
small, so that the proportion of the number of carbides in the
grain boundary cannot be increased. Thus, the cold workability is
low. In the Comparative steel J-2, the annealing temperature during
the second-step annealing during the two-step type annealing is
high, the amount of the carbide remaining is decreased due to the
promoted dissolution of carbides, and pearlite transformation
cannot be suppressed in the structure after the slow cooling, the
Vickers hardness is too high, and the cross-sectional shrinkage
percentage is low. Thus, the cold forgeability is low.
[0188] In the Comparative steel H-2, since the annealing
temperature during the second-step annealing during the two-step
type annealing is low, and the amount of austenite produced is
small, so that the proportion of the number of carbides in the
grain boundary cannot be increased. Thus, the cold workability is
low. In the Comparative steel Y-2, the retention time during the
second-step annealing during the two-step type annealing is long,
the amount of carbides remaining is decreased due to the promoted
dissolution of carbides, and pearlite transformation cannot be
suppressed in the structure after the slow cooling, and the
cross-sectional shrinkage percentage is low. Thus, the cold
forgeability is low. In the Comparative steel E-2, the cooling rate
from the second-step annealing during the two-step type annealing
to 650.degree. C. is fast, pearlite transformation occurred during
cooling, the Vickers hardness is too high, and the cross-sectional
shrinkage percentage is low. Thus, the cold workability is low.
[0189] In any of the comparative steels A-1, D-1, I-1, M-1, A-2 and
I-2, the X-ray diffraction intensity ratio of {311} <011> is
greater than 3.0. In these comparative steels, the in-plane
anisotropy |.DELTA.r| exceeds 0.2, and thus the cold workability is
low. As described above, by performing analysis by X-ray
diffraction on a plane parallel to the plate surface at the 1/2
plate thickness portion of the hot-rolled steel sheet, the degree
of plastic anisotropy such as the in-plane anisotropy |.DELTA.r| or
the quality of cold workability of the hot-rolled steel sheet to be
cold worked can be determined before cold working.
INDUSTRIAL APPLICABILITY
[0190] As described above, according to the present invention, a
steel sheet with excellent cold workability during forming can be
manufactured and provided. The steel sheet of the present invention
is a steel sheet suitable as a material for automotive parts,
blades, and other mechanical parts manufactured through processing
steps such as punching, bending, pressing, etc. Therefore, the
present invention has excellent industrial applicability.
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