U.S. patent application number 11/064049 was filed with the patent office on 2005-09-15 for high carbon hot-rolled steel sheet and method for manufacturing the same.
This patent application is currently assigned to JFE STEEL CORPORATION. Invention is credited to Fujita, Takeshi, Iizuka, Shunji, Matsuoka, Saiji, Nakamura, Nobuyuki, Tsuchiya, Yoshiro.
Application Number | 20050199322 11/064049 |
Document ID | / |
Family ID | 34918380 |
Filed Date | 2005-09-15 |
United States Patent
Application |
20050199322 |
Kind Code |
A1 |
Nakamura, Nobuyuki ; et
al. |
September 15, 2005 |
High carbon hot-rolled steel sheet and method for manufacturing the
same
Abstract
The high carbon hot-rolled steel sheet contains, in terms of
percentages of mass, 0.10 to 0.7% C, 2.0% or less Si, 0.20 to 2.0%
Mn, 0.03% or less P, 0.03% or less S, 0.1% or less Sol.Al, 0.01% or
less N, and the balance being Fe and inevitable impurities, and has
a structure of ferrite having 6 .mu.m or less average grain size
and carbide having 0.10 .mu.m or more and less than 1.2 .mu.m of
average grain size. The volume ratio of the carbide having 2.0
.mu.m or more of grain size is 10% or less. The volume ratio of the
ferrite containing no carbide is 5% or less. The manufacturing
method thereof has the steps of hot-rolling, primary cooling,
holding, coiling, acid washing, and annealing. The primary cooling
step is to cool the hot-rolled steel sheet down to cooling
termination temperatures ranging from 450.degree. C. to 600.degree.
C. at cooling rates of higher than 120.degree. C./sec. The holding
step is to apply secondary cooling to hold the primarily cooled
hot-rolled steel sheet at a temperature range from 450.degree. C.
to 650.degree. C. until coiling.
Inventors: |
Nakamura, Nobuyuki;
(Hiroshima, JP) ; Fujita, Takeshi; (Chiba, JP)
; Tsuchiya, Yoshiro; (Hiroshima, JP) ; Iizuka,
Shunji; (Hiroshima, JP) ; Matsuoka, Saiji;
(Hiroshima, JP) |
Correspondence
Address: |
FRISHAUF, HOLTZ, GOODMAN & CHICK, PC
220 5TH AVE FL 16
NEW YORK
NY
10001-7708
US
|
Assignee: |
JFE STEEL CORPORATION
Tokyo
JP
|
Family ID: |
34918380 |
Appl. No.: |
11/064049 |
Filed: |
February 22, 2005 |
Current U.S.
Class: |
148/636 ;
148/320; 148/654 |
Current CPC
Class: |
C22C 38/22 20130101;
C22C 38/04 20130101; C22C 38/001 20130101; C22C 38/02 20130101;
C22C 38/18 20130101; C22C 38/12 20130101; C22C 38/06 20130101; C21D
8/0226 20130101; C21D 8/0263 20130101 |
Class at
Publication: |
148/636 ;
148/654; 148/320 |
International
Class: |
C21D 008/00 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 10, 2004 |
JP |
2004-067119 |
Claims
1. A high carbon hot-rolled steel sheet consisting essentially of:
in terms of percentage of mass, 0.10 to 0.7% C, 2.0% or less Si,
0.20 to 2.0% Mn, 0.03% or less P, 0.03% or less S, 0.1% or less
Sol.Al, 0.01% or less N, and the balance being Fe and inevitable
impurities; ferrite having an average grain size of 6 .mu.m or less
and carbide having an average grain size of 0.10 .mu.m or more and
less than 1.2 .mu.m; the carbide having a volume ratio of 10% or
less regarding a grain size of 2.0 .mu.m or more; and the ferrite
containing no carbide having a volume ratio of 5% or less.
2. The high carbon hot-rolled steel sheet according to claim 1,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.05 to 1.5% Cr and
0.01 to 0.5% Mo.
3. The high carbon hot-rolled steel sheet according to claim 1,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.005% or less B,
1.0% or less Cu, 1.0% or less Ni, and 0.5% or less W.
4. The high carbon hot-rolled steel sheet according to claim 2,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.005% or less B,
1.0% or less Cu, 1.0% or less Ni, and 0.5% or less W.
5. The high carbon hot-rolled steel sheet according claim 1,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.5% or less Ti,
0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.
6. The high carbon hot-rolled steel sheet according to claim 1,
wherein the content of Si is 0.005 to 2.0% by mass.
7. The high carbon hot-rolled steel sheet according to claim 6,
wherein the content of Si is 0.02 to 0.5% by mass.
8. The high carbon hot-rolled steel sheet according to claim 1,
wherein the content of Mn is 0.2 to 1.0% by mass.
9. The high carbon hot-rolled steel sheet according to claim 2,
wherein the content of Cr is 0.05 to 0.3% by mass.
10. The high carbon hot-rolled steel sheet according to claim 2,
wherein the content of Cr is 0.8 to 1.5% by mass.
11. The high carbon hot-rolled steel sheet according to claim 2,
wherein the content of Mo is from 0.05 to 0.5% by mass.
12. A method for manufacturing a high carbon hot-rolled steel
sheet, comprising the steps of: hot-rolling a steel consisting
essentially of, in terms of percentages of mass, 0.10 to 0.70% C,
2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03% or less S,
0.1% or less Sol.Al, 0.01% or less N, and the balance being Fe and
inevitable impurities, at finishing temperatures of (Ar.sub.3
transformation point -10.degree. C.) or more; applying primary
cooling to the hot-rolled steel sheet down to cooling termination
temperatures ranging from 450.degree. C. to 600.degree. C. at
cooling rates of more than 120.degree. C./sec; applying secondary
cooling to hold the primarily cooled hot-rolled steel sheet in a
temperature range from 450.degree. C. to 650.degree. C. until
coiling; coiling the cooled hot-rolled steel sheet at coiling
temperatures of 600.degree. C. or less; applying acid washing to
the coiled hot-rolled steel sheet; and annealing the acid-washed
hot-rolled steel sheet at annealing temperatures ranging from
680.degree. C. to Ac.sub.1 transformation point.
13. The method according claim 12, wherein the cooling rate in the
primary cooling step is in a range from 120 to 700.degree.
C./sec.
14. The method according to claim 12, wherein the coiling
temperature is in a range from 200.degree. C. to 600.degree. C.
15. The high carbon hot-rolled steel sheet according claim 2,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.5% or less Ti,
0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.
16. The high carbon hot-rolled steel sheet according claim 3,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.5% or less Ti,
0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.
17. The high carbon hot-rolled steel sheet according claim 4,
further containing at least one element selected from the group
consisting of, in terms of percentages of mass, 0.5% or less Ti,
0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.
Description
FIELD OF THE INVENTION
[0001] The present invention relates to a high carbon hot-rolled
steel sheet having excellent ductility and stretch-flange
formability, and a manufacturing method thereof.
DESCRIPTION OF THE RELATED ARTS
[0002] High carbon steel sheets employed for tools, automobile
parts (gear, transmission), and the like are subjected to heat
treatment such as quenching and tempering after punching and
forming thereof. The requests of users who conduct the working on
these components include improvement in bore expansion (burring)
property in the forming process after punching, as well as the
elongation characteristic which is an index of ductility for
forming the steel sheet into complex shapes. The burring property
is evaluated by the stretch-flange formability as one of
press-forming properties. Consequently, there are wanted the
materials having excellent stretch-flange formability as well as
ductility.
[0003] Regarding the improvement in the stretch-flange formability
of high carbon steel sheets, several technologies have been
studied. For example, JP-A-11-269552 and JP-A-11-269553, (the term
"JP-A" referred to herein signifies the "Japanese Patent Laid-Open
Publication"), disclose a method for manufacturing medium to high
carbon steel sheets having excellent stretch-flange formability in
a process after cold-rolling. The disclosed technology employs a
hot-rolled steel which contains 0.1 to 0.8% C by mass, having a
metallic structure substantially consisting of ferrite phase and
pearlite phase, having, at need, the area rate of proeutectoid
ferrite of at or higher value determined by the C content (% by
mass), and having 0.1 .mu.m or larger distance between pearlite
lamellas. To the hot-rolled steel sheet, cold-rolling is given by
15% or higher rolling rate, followed by three-stage or two-stage
annealing while holding the steel sheet in three steps or two steps
of temperature ranges for a long time.
[0004] Also JP-A-2003-13145 discloses a method for manufacturing a
high carbon steel sheet having excellent stretch-flange
formability, which contains 0.2 to 0.7% C by mass, has average
grain size of carbide in a range from 0.1 to 1.2 .mu.m, and has a
volume ratio of carbide-free ferrite grains of 10% or less. The
disclosed technology is a process in which the hot-rolling is given
at finishing temperatures of (Ar.sub.3 transformation point
-20.degree. C.) or above, the cooling is given to cooling
termination temperatures of 650.degree. C. or below at cooling
rates of higher than 120.degree. C./sec, the coiling is given at
temperatures of 600.degree. C. or below, the acid washing is given,
and then the annealing is given at annealing temperatures ranging
from 640.degree. C. to Ac.sub.1 transformation point.
[0005] According to the technologies disclosed in JP-A-11-269552
and JP-A-11-269553, the ferrite structure is made by the
proeutectoid ferrite and does not include carbide. As a result, the
stretch-flange formability is not necessarily favorable, though the
material is soft and shows excellent ductility. A presumable reason
of the phenomenon is the following. During punching the steel
sheet, the area of proeutectoid ferrite significantly deforms in
the vicinity of a punched end face, which induces significant
difference between the deformation of the proeutectoid ferrite and
that of the ferrite containing spheroid carbide. As a result,
stress concentrates to the peripheral zones of grain boundary where
the deformation significantly differs therebetween, thereby
generating voids at interface between the spheroid structure and
the ferrite. Since the voids grow to cracks, the stretch-flange
formability is ultimately deteriorated.
[0006] A countermeasure to the phenomenon may be the one to apply
strengthened spheroidizing annealing, thereby softening the entire
structure. In this measure, however, the spheroidized carbide
becomes coarse to become the origin of void during the forming
step, and the carbide becomes less soluble in the heat treatment
step after the forming to cause the decrease in quenched
strength.
[0007] Furthermore, recent requirement for the forming level has
increased more than ever from the point of increase in
productivity. Consequently, burring in high carbon steel sheet also
likely induces crack generation at punched end face caused by the
advanced level of forming. Therefore, the high carbon steel sheets
are also requested to have high stretch-flange formability.
[0008] In this regard, the inventors of the present invention
developed a technology disclosed in JP-A-2003-13145 aiming to
provide a high carbon steel sheet having excellent stretch-flange
formability and inducing very few cracks at punched end face, which
steel sheet is manufactured without applying time-consuming
multi-stage annealing. The technology allowed manufacturing a high
carbon hot-rolled steel sheet having excellent stretch-flange
formability.
[0009] On the other hand, recent uses of driving system components
and the like request increased strength also in the non-heat
treating parts, specifically in integrally formed components for
attaining higher durability and lighter weight, thus the steel
sheets as the starting material are requested to have 440 MPa or
higher tensile strength (TS). That kind of request with the aim to
reduce the manufacturing cost of components has led a request to
supply hot-rolled steel sheets.
[0010] The integral forming process has more than ten pressing
steps, and is conducted in a complex combination of forming modes
including not only burring but also stretching and bending.
Accordingly, the integral forming has faced the simultaneous
requests of stretch-flangeability and elongation.
[0011] According to the technology disclosed in JP-A-2003-13145,
however, achieving TS=440 MPa (73 point or more as HRB hardness)
not necessarily attains satisfactory stretch-flange formability.
That is, the technology cannot satisfy stably the requirements of
both that level of TS and the stretch-flange formability.
Furthermore, the disclosed technology does not refer to the
elongation.
[0012] Adding to the above technology, the technology disclosed in
JP-A-2003-13145 generates transformation heat after cooling, which
increases the temperature to enhance the precipitation of
proeutectoid ferrite and the pearlite transformation, thereby
inducing growth of coarse carbide and uneven carbide distribution
to likely deteriorate the characteristics.
SUMMARY OF THE INVENTION
[0013] It is an object of the present invention to provide a high
carbon hot-rolled steel sheet having 440 MPa or higher tensile
strength and giving excellent ductility and stretch-flange
formability, generating very few cracks at punched end face, and
which steel sheet can be manufactured without applying
time-consuming multi-stage annealing.
[0014] The inventors of the present invention conduced intensive
studies on the effect of components and microscopic structures of
high carbon steel sheet on ductility and stretch-flange formability
while securing strength thereof, and found that the ductility and
the stretch-flange formability of steel sheet are significantly
affected by not only the composition of the steel, the shape and
quantity of carbide, but also the dispersed state of carbide. That
is, it was found that the ductility and the stretch-flange
formability of high carbon hot-rolled steel sheet are improved by
controlling each of the carbide shape in terms of average grain
size of carbide and volume ratio of carbide having 2.0 .mu.m or
larger grain size, and the dispersed state of carbide in terms of
volume ratio of carbide-free ferrite grains and average grain size
of ferrite.
[0015] The present invention provides a high carbon hot-rolled
steel sheet consisting essentially of, in terms of percentages of
mass, 0.10 to 0.7% C, 2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or
less P, 0.03% or less S, 0.1% or less Sol.Al, 0.01% or less N, and
balance of Fe and inevitable impurities, and having a structure of
ferrite having 6 .mu.m or smaller average grain size and carbide
having 0.10 .mu.m or larger and smaller than 1.2 g m of average
grain size. The volume ratio of the carbide having 2.0 .mu.m or
larger grain size is 10% or less, and the volume ratio of the
ferrite containing no carbide is 5% or less. The high carbon steel
sheet gives excellent ductility and stretch-flange formability.
[0016] The high carbon hot-rolled steel sheet may further contain
at least one element selected from the group consisting of, in
terms of percentages of mass, 0.05 to 1.5% Cr and 0.01 to 0.5%
Mo.
[0017] The high carbon hot-rolled steel sheet may further contain
at least one element selected from the group consisting of, in
terms of percentages of mass, 0.005% or less B, 1.0% or less Cu,
1.0% or less Ni, and 0.5% or less W.
[0018] The high carbon hot-rolled steel sheet may further contain
at least one element selected from the group consisting of, in
terms of percentages of mass, 0.05 to 1.5% Cr and 0.01 to 0.5% Mo,
and further at least one element selected from the group consisting
of, in terms of percentages of mass, 0.005% or less B, 1.0% or less
Cu, 1.0% or less Ni, and 0.5% or less W.
[0019] The high carbon hot-rolled steel sheet may further contain
at least one element selected from the group consisting of, in
terms of percentages of mass, 0.5% or less Ti, 0.5% or less Nb,
0.5% or less V, and 0.5% or less Zr.
[0020] The content of Si is preferably from 0.005 to 2.0% by mass.
From the point of securing strength after annealing, the Si content
is more preferably 0.02% or more. From the point of surface
property, the Si content is more preferably 0.5% or less.
[0021] The content of Mn is preferably from 0.2 to 1.0% by
mass.
[0022] A preferable range of the content of Cr is determined from
the viewpoint of securing sufficient strength after quenching.
Under a condition of securing satisfactory cooling rate in
quenching treatment, the content of Cr is preferably from 0.05 to
0.3% by mass. When the strength after quenching is strictly
required even under varied cooling rate in the quenching treatment,
the Cr content is preferably from 0.8 to 1.5% by mass.
[0023] The content of Mo is preferably from 0.05 to 0.5% by
mass.
[0024] The present invention further provides a method for
manufacturing high carbon hot-rolled steel sheet, having the steps
of hot-rolling, primary cooling, holding, coiling, acid washing,
and annealing.
[0025] The hot-rolling step applies hot-rolling to a steel
consisting essentially of, in terms of percentages of mass, 0.10 to
0.70% C, 2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03%
or less S, 0.1% or less Sol.Al, 0.01% or less N, and balance of Fe
and inevitable impurities, at finishing temperatures of (Ar.sub.3
transformation point -10.degree. C.) or above.
[0026] The steel may further contain at least one element selected
from the group consisting of, in terms of percentages of mass, 0.05
to 1.5% Cr and 0.01 to 0.5% Mo.
[0027] The steel may further contain at least one element selected
from the group consisting of, in terms of percentages of mass,
0.005% or less B, 1.0% or less Cu, 1.0% or less Ni, and 0.5% or
less W.
[0028] The steel may further contain at least one element selected
from the group consisting of, in terms of percentages of mass, 0.05
to 1.5% Cr and 0.01 to 0.5% Mo, and further at least one element
selected from the group consisting of, in terms of percentages of
mass, 0.005% or less B, 1.0% or less Cu, 1.0% or less Ni, and 0.5%
or less W.
[0029] The steel may further contain at least one element selected
from the group consisting of, in terms of percentages of mass, 0.5%
or less Ti, 0.5% or less Nb, 0.5% or less V, and 0.5% or less
Zr.
[0030] The primary cooling step is primary cooling of a hot-rolled
steel sheet down to the cooling termination temperatures ranging
from 450.degree. C. to 600.degree. C. at cooling rates of higher
than 120.degree. C./sec. The upper limit of the cooling rate is
preferably 700.degree. C./sec from the point of facility
capacity.
[0031] The holding step is to hold the cooled hot-rolled steel
sheet in a temperature range from 450.degree. C. to 650.degree. C.
by the secondary cooling until coiling.
[0032] The coiling step is to coil the cooled hot-rolled steel
sheet at coiling temperatures of 600.degree. C. or below. The
coiling temperature is preferably in a range from 200.degree. C. to
600.degree. C.
[0033] The acid washing step is to apply acid washing to the coiled
hot-rolled steel sheet.
[0034] The annealing step is to anneal the hot-rolled steel sheet
after the acid washing at temperatures ranging from 680.degree. C.
to Ac.sub.1 transformation point.
[0035] The percentage indicating the composition of steel, referred
to herein, is percentage by mass.
[0036] The present invention suppresses the generation of voids at
punched end face during punching, and delays the growth of cracks
during burring. As a result, the present invention provides a high
carbon hot-rolled steel sheet having 440 MPa or higher tensile
strength and extremely excellent ductility and stretch-flange
formability. By applying the high carbon hot-rolled steel sheet
having excellent ductility and stretch-flange formability according
to the present invention to highly durable parts such as
transmission parts represented by gear, advanced level of forming
is attained in the forming step, which provides high product
quality and allows manufacturing the parts at low cost with
decreased number of manufacturing steps. Also for the parts of
driving system, the integrally formed components are requested to
have increased strength in the non-heat treating parts for
attaining higher durability and lighter weight, thus the steel
sheets as the starting material are requested to have 440 MPa class
tensile strength (TS). The high carbon hot-rolled steel sheet
according to the present invention is useful in this respect.
DESCRIPTION OF THE EMBODIMENTS
[0037] The high carbon hot-rolled steel sheet according to the
present invention consists essentially of, in terms of percentages
of mass, 0.1 to 0.7% C, 2.0% or less Si, 0.2 to 2.0% Mn, 0.03% or
less P, 0.03% or less S, 0.1% or less Sol.Al, 0.01% or less N, and
balance of Fe and inevitable impurities, and has a structure of
ferrite having 6 .mu.m or smaller average grain size and carbide
having 0.10 .mu.m or more and less than 1.2 .mu.m of average grain
size, wherein the volume ratio of the carbide having 2.0 .mu.m or
larger grain size is 10% or less, and the volume ratio of the
ferrite containing no carbide is 5% or less. The above
specification of the steel sheet is most important parameter of the
present invention. With thus specified chemical composition,
metallic structure (average grain size of ferrite), shape of
carbide (volume ratio of carbide having 2.0 .mu.m or larger average
grain size), and dispersion state of carbide (volume ratio of
carbide-free ferrite grains), and by satisfying all of these
specifications, a high carbon hot-rolled steel sheet having
excellent ductility and stretch-flange formability is obtained.
[0038] The high carbon hot-rolled steel sheet according to the
present invention may further contain one or both of, in terms of
percentage by mass, 0.05 to 1.5% C and 0.01 to 0.5% Mo, may further
contain one or more of, in terms of percentage by mass, 0.005% or
less B, 1.0% or less Cu, 1.0% or less Ni, and 0.5% or less W, and
may further contain one or more of, in terms of percentage by mass,
0.5% or less Ti, 0.5% or less Nb, 0.5% or less V, and 0.5% or less
Zr.
[0039] The high carbon hot-rolled steel sheet can be manufactured
by the steps of: hot-rolling the steel at finishing temperatures of
(Ar.sub.3 transformation point -10.degree. C.) or above; applying
primary cooling to the hot-rolled steel sheet down to cooling
termination temperatures ranging from 450.degree. C. to 600.degree.
C. at cooling rates of higher than 120.degree. C./sec; applying
secondary cooling to hold the primarily cooled hot-rolled steel
sheet in a temperature range from 450.degree. C. to 650.degree. C.
until coiling; coiling the cooled hot-rolled steel sheet at coiling
temperatures of 600.degree. C. or below; applying acid washing to
the coiled hot-rolled steel sheet; and annealing the acid-washed
hot-rolled steel sheet at annealing temperatures ranging from
680.degree. C. to Ac.sub.1 transformation point. The object of the
invention is attained by totally controlling the conditions of,
after the hot-rolling, primary cooling, secondary cooling, coiling,
and annealing.
[0040] The present invention is described in more detail in the
following.
[0041] The reasons to limit the chemical composition of the steel
according to the present invention are described below.
[0042] C: 0.1 to 0.7%
[0043] Carbon is an important element that forms carbide and
provides hardness after quenching. However, the C content of less
than 0.1% causes conspicuous formation of proeutectoid ferrite in
the structure after the hot-rolling, which results in uneven
carbide distribution. In such a case, strength sufficient for
structural machine parts cannot be obtained even after quenching.
On the other hand, the C content exceeding 0.7% results in
insufficient working property, giving low stretch-flange
formability and ductility. In such a case, the steel sheet after
the hot-rolling shows high hardness and becomes brittle so that the
strength after quenching saturates. Therefore, the C content is
specified to a range from 0.1 to 0.7%. From the point of securing
sufficient strength after quenching, the C content is preferably
0.2,% or more, and from the point of handling of steel sheet on and
after coiling, the C content is preferably 0.6% or less. The C
content condition is an important parameter of the present
invention.
[0044] Si: 2.0% or Less
[0045] Since Si is an element to improve the quenching property and
increase the material strength by solid solution strengthening, the
Si content is preferably 0.005% or more. However, The Si content
exceeding 2.0% facilitates formation of proeutectoid ferrite and
increases the ferrite grains substantially free from carbide,
thereby deteriorating the stretch-flange formability. Furthermore,
Si has a tendency of graphitizing carbide and likely hinders
quenching property. Consequently, the Si content is specified to
2.0% or less, preferably 0.02% or more from the point of securing
strength after annealing, and preferably 0.5% or less from the
point of surface property.
[0046] Mn: 0.2 to 2.0%
[0047] Similar with Si, Mn is an element to improve the quenching
property and to increase the material strength by solid solution
strengthening. Manganese is also an important element which fixes S
as MnS and prevents hot cracking of slab. However, the Mn content
of less than 0.2% reduces these effects, and enhances the formation
of proeutectoid ferrite to generate coarse ferrite grains, and
further significantly deteriorates the quenching property. The Mn
content exceeding 2.0% allows significant formation of manganese
band which is a segregation zone, though a wanted tensile strength
is attained, thereby deteriorating the stretch-flange formability
and the elongation. Accordingly, the Mn content is specified to a
range from 0.20% to 2.0%, and preferably 1.0% or less from the
viewpoint of stretch-flange formability and deterioration in
elongation caused by the formation of manganese band.
[0048] P: 0.03% or Less
[0049] Phosphorus is an element to be reduced because P is
segregated in grain boundaries to decrease the toughness. Since,
however, the P content is acceptable up to 0.03%, the P content is
specified to 0.03% or less.
[0050] S: 0.03% or Less
[0051] Sulfur is an element to be reduced because S forms MnS with
Mn to deteriorate the stretch-flange formability. Since, however,
the S content is acceptable up to 0.03%, the S content is specified
to 0.03% or less.
[0052] sol.Al: 0.1% or less
[0053] Aluminum is added in the steel making stage as an
acid-eliminating agent to improve the cleanliness of steel.
Normally Al is contained in the steel in an amount of 0.005% or
more as sol.Al. An Al content exceeding 0.1% as sol.Al results in
the saturation of the cleanliness improving effect, thereby
increasing the cost. In addition, excess Al results in large amount
of AlN precipitate to deteriorate the quenching property.
Therefore, the sol.Al content is specified to 0.1% or less,
preferably 0.08% or less.
[0054] N: 0.01% or Less
[0055] Since excess N deteriorates the ductility, the N addition is
specified to 0.01% or less.
[0056] The steel sheet according to the present invention achieves
the objective characteristics with the above essential adding
elements. Depending on the wanted characteristics, however, one or
both of Cr and Mo may be added.
[0057] Cr: 0.05 to 1.5%
[0058] Chromium is an important element to suppress the formation
of proeutectoid ferrite during cooling step after the hot-rolling,
thus to improve the stretch-flange formability and improve the
quenching property. However, the Cr content less than 0.05% cannot
attain satisfactory effect. Furthermore, the Cr content exceeding
1.5% saturates the effect to suppress the formation of proeutectoid
ferrite and increases the cost, though the quenching property is
improved. Accordingly, when Cr is added, the Cr content is
specified to a range from 0.05 to 1.5%. Preferably, from the point
of securing sufficient strength after quenching, the Cr content is
in a range from 0.05 to 0.3% under a condition that a satisfactory
cooling rate is assured at quenching, and from 0.8 to 1.5% when a
strict strength condition is requested after quenching even under
varied cooling rate at quenching.
[0059] Mo: 0.01 to 0.5%
[0060] Molybdenum is an important element to suppress the formation
of proeutectoid ferrite during the cooling step after the
hot-rolling, thus to improve the stretch-flange formability and
improve the quenching property. However, the Mo content of less
than 0.01% cannot attain satisfactory effect. On the other hand,
the Mo content exceeding 0.5% saturates the effect to suppress the
formation of proeutectoid ferrite and increases the cost, though
the quenching property is improved. Accordingly, when Mo is added,
the Mo content is specified to a range from 0.01 to 0.5%, and
preferably 0.05% or more from the point of securing sufficient
strength after quenching.
[0061] The steel according to the present invention may further
contain, adding to the above adding elements, one or more of B, Cu,
Ni, and W, at need, to suppress the formation of proeutectoid
ferrite during hot-rolling and cooling and to improve the quenching
property. In such a case, less than 0.0001% B, and less than 0.01%
for each of Cu, Ni, and W cannot fully attain the added effect. On
the other hand, the added quantity exceeding 0.005% B, 1.0% Cu,
1.0% Ni, and 0.5% W saturates the added affect, and increases the
cost. Consequently, on adding these elements, the specified content
is 0.0001 to 0.005% B, 0.01 to 1.0% Cu, 0.01 to 1.0% Ni, and 0.01
to 0.5% W. Boron, however, may form a compound with N in the steel
to fail in providing the effect of B itself. Therefore, the element
to be added for suppressing the formation of proeutectoid ferrite
during hot-rolling and cooling and for improving the quenching
property is preferably selected by one or more among the elements
of Cu, Ni, and W. In that case, preferable adding amount of the
respective elements is similar with that given above.
[0062] The steel according to the present invention may further
contain, adding to the above adding elements, one or more of Ti,
Nb, V, and Zr for assuring 440 MPa or higher tensile strength by
refining the ferrite grains. In that case, each content less than
0.001% cannot obtain sufficient effect of addition. On the other
hand, each content exceeding 0.5% saturates the adding effect and
increases the cost. Therefore, if these elements are added, the
content of each one is specified to a range from 0.001 to 0.5%.
[0063] The balance to the above composition is Fe and inevitable
impurities.
[0064] During the manufacturing process, various elements such as
Sn and Pb may enter as impurities. Those kinds of impurities,
however, do not influence the effect of the present invention.
[0065] The following is the description of the present invention in
terms of metallic structure (average grain size of ferrite), shape
of carbide (average grain size of carbide and volume ratio of
carbide having 2.0 .mu.m or larger average grain size), and
dispersion state of carbide (volume ratio of carbide-free ferrite
grains). These conditions are important parameters to obtain the
high carbon hot-rolled steel sheet having excellent ductility and
stretch-flange formability, and the effect of the present invention
cannot be attained if any of these conditions is not satisfied, or
the effect of the present invention is attained only after
satisfying all of these conditions.
[0066] Average Ferrite Grain Size: 6 .mu.m or Smaller
[0067] The average ferrite grain size is an important parameter
governing the stretch-flange formability and the material strength.
By refining the ferrite grains, the strength is increased without
deteriorating the stretch-flange formability. More specifically,
average ferrite grain sizes of 6 .mu.m or smaller provide excellent
ductility and stretch-flange formability while securing 440 MPa or
higher tensile strength of the material. The average ferrite grain
size can be controlled by the primary cooling termination
temperature, the secondary cooling holding temperature, and the
coiling temperature, after hot-rolling, which are described
below.
[0068] Average Carbide Grain Size: 0.10 .mu.m or Larger and Smaller
Than 1.2 .mu.m
[0069] The average carbide grain size significantly influences the
working properties in general and the void formation during
burring. Thus the average carbide grain size is an important
parameter of the present invention. Although smaller carbide grain
sizes suppress more the void formation, average carbide grain size
of smaller than 0.10 .mu.m deteriorates the ductility with the
increase in hardness, thereby deteriorating the stretch-flange
formability. On the other hand, increased average carbide grain
size generally improves the working property. The size exceeding
1.2 .mu.m, however, leads to void formation during burring to
deteriorate the stretch-flange formability, and further the
decrease in the local ductility causes the deterioration of
ductility. Consequently, the average carbide grain size is
specified to a range from 0.10 .mu.m or larger and smaller than 1.2
.mu.m. As described below, the average carbide grain size can be
controlled by the manufacturing conditions, specifically by the
primary cooling termination temperature, the coiling temperature,
and the annealing temperature.
[0070] Volume Ratio of Carbide Having 2.0 .mu.m or Larger Grain
Size: 10% or Less
[0071] During general working process and burring step, voids
predominantly occur in the vicinity of coarse carbide. Accordingly,
carbide has to be emphasized to control the average grain size and
to reduce the volume ratio of coarse carbide grains, and they are
also important parameters of the present invention. Even when the
average carbide grain size is in a range from 0.10 .mu.m or larger
and smaller than 1.2 .mu.m, the existence of more than 10% volume
ratio of coarse carbide grains at or larger than 2.0 .mu.min size
deteriorates the stretch-flange formability caused by the
generation of voids during burring, thereby decreasing the local
ductility to result in the deterioration of ductility.
Consequently, the volume ratio of the carbide having 2.0 .mu.m or
larger grain size is specified to 10% or less. As described below,
the carbide grain size can be controlled by the primary cooling
termination temperature, the secondary cooling holding temperature,
the coiling temperature, and the annealing temperature.
[0072] Volume Ratio of Carbide-Free Ferrite Grain Size: 5% or
Less
[0073] Uniform dispersion of carbide relaxes the stress
concentration on a punched end face during burring, thereby
suppressing the void formation. In this regard, it is important to
control the volume ratio of carbide-free ferrite grains. By
controlling the volume ratio of carbide-free ferrite grains to 5%
or less, the effect similar with the state of uniform dispersion of
carbide is attained, and the stretch-flange formability is
significantly improved. In addition, local ductility is improved,
which then significantly improves the ductility. The term
"carbide-free" referred to herein signifies that no carbide is
detected in an ordinary metal structure observation (with an
optical microscope). That type of ferrite grains forms a zone
appeared as the proeutectoid ferrite after hot-rolling, where
substantially no carbide is observed within grain even after the
annealing. As described below, the state of carbide dispersion can
be controlled by the manufacturing conditions, specifically by the
finishing temperature, the cooling rate during cooling after the
rolling, the cooling termination temperature, and the coiling
temperature.
[0074] The following is the description about the manufacturing
method for high carbon hot-rolled steel sheet having excellent
ductility and stretch-flange formability according to the present
invention.
[0075] The high carbon hot-rolled steel sheet according to the
present invention is obtained by the steps of: hot-rolling a steel
prepared to have the above range of chemical composition at
finishing temperatures of (Ar.sub.3 transformation point
-10.degree. C.) or above; applying primary cooling to the
hot-rolled steel sheet down to cooling termination temperatures
ranging from 450.degree. C. to 600.degree. C. at cooling rates of
higher than 120.degree. C./sec; applying secondary cooling to hold
the primarily cooled hot-rolled steel sheet in a temperature range
from 450.degree. C. to 650.degree. C. until coiling; coiling the
cooled hot-rolled steel sheet at coiling temperatures of
600.degree. C. or below; applying acid washing to the coiled
hot-rolled steel sheet; and annealing the acid-washed hot-rolled
steel sheet at annealing temperatures ranging from 680.degree. C.
to Ac.sub.1 transformation point. The detail of the respective
steps is described below.
[0076] Finishing Temperature: Hot-Rolling at (Ar.sub.3
Transformation Point -10.degree. C.) or Above
[0077] Finishing temperature of hot-rolling below (Ar.sub.3
transformation point -10.degree. C.) enhances the ferrite
transformation in a part, which increases the ferrite grains to
deteriorate the ductility and the stretch-flange formability.
Therefore, the finish-rolling is done at finishing temperatures of
(Ar.sub.3 transformation point -10.degree. C.) or above. The
condition assures uniform structure and improves the ductility and
the stretch-flange formability.
[0078] Cooling Rate: Primary Cooling at Rates of Higher Than
120.degree. C./sec
[0079] According to the present invention, rapid cooling (primary
cooling) is adopted at cooling rates of higher than 120.degree.
C./sec after hot-rolling to reduce the volume ratio of proeutectoid
ferrite after transformation. Gradual cooling results in a low
super cooling degree of austenite, leading to the formation of
proeutectoid ferrite. In particular, 0.120.degree. C./sec or
smaller cooling rate gives conspicuous formation of proeutectoid
ferrite, thereby resulting in the carbide-free ferrite grains
exceeding 5% to deteriorate the ductility and the stretch-flange
formability. Accordingly, the cooling rate after hot-rolling is
specified to higher than 120.degree. C./sec.
[0080] It is preferable to begin the primary cooling after the
finish-rolling within a period of from more than 0.1 sec and less
than 1.0 sec. The condition provides finer ferrite grains and
precipitates such as pearlite after the transformation, thus
further improving the working property.
[0081] Cooling Termination Temperature: 450.degree. C. to
600.degree. C.
[0082] High cooling termination temperature in the primary cooling
causes proeutectoid ferrite formation and increase in the lamella
spacing of pearlite. As a result, fine carbide cannot be obtained
after the annealing, and the ductility and the stretch-flange
formability are deteriorated. Particularly when the cooling
termination temperature is higher than 600.degree. C., the
carbide-free ferrite grains increase to more than 5%, which
deteriorates the ductility and the stretch-flange formability.
Therefore, the cooling termination temperature after rolling is
specified to 600.degree. C. or below. Lower than 450.degree. C. of
cooling termination temperature cannot obtain the equiaxed ferrite
grains, and deteriorates the working property. Therefore, the
cooling termination temperature is specified to 450.degree. C. or
above.
[0083] Secondary Cooling From the Primary Cooling Termination to
the Coiling: Holding at Temperatures in a Range From 450.degree. C.
to 650.degree. C.
[0084] For the case of high carbon steel sheets, the steel sheet
temperature increases after the primary cooling termination, in
some cases, accompanied by the proeutectoid ferrite transformation,
the pearlite transformation, and the bainite transformation. Thus,
even if the primary cooling termination temperature is lower than
600.degree. C., when the temperature in the course from the primary
cooling termination to the coiling is higher than 650.degree. C.,
the proeutectoid ferrite is formed, the lamella spacing of pearlite
increases, and the carbide in pearlite becomes coarse. As a result,
the fine carbide cannot be obtained after the annealing, and the
volume ratio of carbide having 2.0 .mu.m or larger grain size
exceeds 10%, thereby deteriorating the ductility and the
stretch-flange formability. If the temperature in the course from
the primary cooling termination to the coiling is lower than
450.degree. C., the equiaxed ferrite cannot be obtained to
deteriorate the working property, in some cases. Therefore, it is
important to control the temperature in the course from the
secondary cooling to the coiling. By holding the material between
the secondary cooling step and the coiling step to temperatures
ranging from 450.degree. C. to 650.degree. C., the deterioration of
ductility, of stretch-flange formability, and of working property
can be prevented. The secondary cooling may be done by laminar
cooling or the like.
[0085] Regarding the holding time from the primary cooling
termination to the coiling, short in the time induces the
generation of transformation heat after coiling, which makes the
steel sheet temperature control impossible and generates coil
crushing. Therefore, the holding time is preferably 5 seconds or
more for completing the transformation until coiling, and
preferably 60 seconds or less because excess holding time
significantly deteriorates the operability.
[0086] Coiling Temperature: 600.degree. C. or below
[0087] Higher coiling temperature increases more the lamella
spacing of pearlite. Thus, the carbide becomes coarse after the
annealing. When the coiling temperature exceeds 600.degree. C., the
ductility and the stretch-flange formability deteriorate.
Consequently, the coiling temperature is specified to 600.degree.
C. or below. Although the lower limit of the coiling temperature is
not specifically defined, 200.degree. C. or above is preferred
because lower temperature induces more the deterioration of steel
sheet shape.
[0088] Annealing Temperature: 680.degree. C. to Ac.sub.1
Transformation Point
[0089] After applying acid washing to the hot-rolled steel sheet,
annealing is given for spheroidizing the carbide. The annealing
temperature lower than 680.degree. C. results in insufficient
spheroidization of carbide or in forming carbide having smaller
than 0.1 .mu.m of average grain size, which deteriorates the
stretch-flange formability. In addition, no equiaxed ferrite is
obtained, and the working property and the ductility are
deteriorated. On the other hand, annealing temperature exceeding
the Ac.sub.1 transformation point causes austenite formation in a
part, which again generates pearlite during cooling, thereby also
deteriorating the stretch-flange formability and the ductility.
Consequently, the annealing temperature is specified to a range
from 680.degree. C. to Ac.sub.1 transformation point.
[0090] For the composition preparation of the high carbon steel
according to the present invention, either a converter or an
electric furnace can be applied. The high carbon steel after the
composition preparation is formed in a steel slab by block
formation--block rolling or by continuous casting. The steel slab
is subjected to hot-rolling. The slab heating temperature is
preferably 1280.degree. C. or below to avoid deterioration of the
surface state caused by scaling. The continuously cast slab may be
sent, in as-cast state, to direct-feed rolling in which the slab is
rolled under heating to prevent temperature reduction. Furthermore,
finish-rolling may be given during the hot-rolling step eliminating
the rough-rolling. Alternatively, to secure the finishing
temperature, the rolled material may be heated with a heating means
such as bar heater during the hot-rolling. Also in order to
accelerate spheroidization or to reduce the hardness, the coiled
steel sheet may be held to the temperature with a gradual cooling
cover or other means.
[0091] The annealing after hot-rolling may be conducted by box
annealing or continuous annealing. Temper rolling is succeedingly
executed at need. Since the temper rolling does not influence the
quenching property, the condition of temper rolling is not
specifically limited.
[0092] The above procedure provides a high carbon hot-rolled steel
sheet having excellent ductility and stretch-flange formability. A
presumable reason that the high carbon hot-rolled steel sheet
according to the present invention has the excellent ductility and
stretch-flange formability is the following. The stretch-flange
formability is significantly affected by the internal structure of
punched end face zone. It was confirmed that, particularly for the
case of large amount of carbide-free ferrite grains (the
proeutectoid ferrite after the hot-rolling), cracks are generated
from the grain boundary with the spheroidal structure zone. When
the behavior of microstructure is observed, the void formation
caused by the stress concentration becomes stronger at the
interface of carbide after the punching. The stress concentration
is enhanced in a state of increased size of carbide grains and
increased quantity of carbide-free ferrite grains. On burring,
these voids are connected each other to form cracks. Further by
controlling the ferrite grain size, the elongation stably
increases. From the above phenomena, it is possible to reduce
stress concentration, to reduce void generation, thus to provide
excellent ductility and stretch-flange formability through the
control of chemical composition, metallic structure (average
ferrite grain size), carbide shape (volume ratio of carbide having
2.0 .mu.m or larger average grain size), and dispersed state of
carbide (volume ratio of carbide-free ferrite grains).
EXAMPLE 1
[0093] Continuously cast slabs of steels having the respective
chemical compositions given in Table 1 as the steel Nos. A to R
were heated to 1250.degree. C., then were subjected to hot-rolling
and annealing under the respective conditions given in Table 2 to
prepare steel sheets having 5.0 mm in thickness. The steel sheet
Nos. 1 to 18 are the example steels prepared under the
manufacturing conditions within the range of the present invention,
and the steel Nos. 19 to 32 are the comparative example steels
prepared under the manufacturing conditions outside the range of
the present invention.
[0094] Samples were cut from thus prepared respective steel sheets,
and were subjected to measurements of ferrite grain size, average
carbide grain size, volume ratio of carbide having 2.0 2 m or
larger grain size, volume ratio of carbide-free ferrite grains,
hardness, and stretch-flange formability (burring ratio), and
further to tensile test. The results are given in Table 3. Method
and condition of each test and measurement are the following.
[0095] (1) Determination of Ferrite Grain Size, Average Carbide
Grain Size, Volume Ratio of Carbide Having 2.0 .mu.m or Larger
Grain Size, and Volume Ratio of Carbide-Free Ferrite Grains
[0096] A cross section along the thickness of a sample sheet was
polished, etched, and photographed by a scanning electron
microscope to observe the microstructure within an area of 0.01
mm.sup.2; The determination was given on the ferrite grain size,
the average carbide grain size, the volume ratio of carbide having
2.0 .mu.m or larger grain size, and the volume ratio of
carbide-free ferrite grains.
[0097] (2) Determination of Hardness
[0098] The surface hardness of steel sheet was determined in
accordance with JIS Z2245. Average of n=5 data was derived.
[0099] (3) Determination of Stretch-Flange Formability
[0100] A sample was punched with a punching tool having a punch
diameter of d.sub.0=10 mm and a die diameter of 12 mm (clearance
20%), and was subjected to a hole-expanding test. The
hole-expanding test was executed by the push-up method with a
cylindrical flat-bottomed punch (50 mmf, 8R)), then a hole diameter
db was measured when a crack was generated across the thickness of
the sheet. The hole-expanding ratio .lambda.(%) defined by the
following formula was derived.
.lambda.=100.times.(db-d.sub.0)/d.sub.0 (1)
[0101] (4) Tensile Test
[0102] A JIS No. 5 sheet was cut along the direction of 90.degree.
(C direction) to the rolling direction, and was subjected to
tensile test with a testing speed of 10 mm/min to determine the
tensile strength and the elongation.
[0103] The present invention places the target values of: 440 MPa
or higher tensile strength TS; 35% or higher elongation for a steel
containing 0.10% or more and less than 0.40% C; 30% or higher
elongation for a steel containing 0.40 to 0.70% C; 70% or higher
hole-expanding ratio .lambda. for a steel containing 0.10% or more
and less than 0.40% C (5.0 mm of sheet thickness); and 40% or
higher hole expanding ratio .lambda. for a steel containing 0.40 to
0.70% C (5.0 mm of sheet thickness).
[0104] Table 3 shows that the example steel sheet Nos. 1 to 18 of
the present invention gave 440 MPa or higher tensile strength (TS),
with high hole-expanding ratio .lambda., thus providing excellent
stretch-flange formability and elongation.
[0105] In contrast, the steel sheet Nos. 19 to 32 are the
comparative example steels which were prepared under the
manufacturing conditions outside the range of the present
invention. The steel sheet Nos. 19, 20, 22, 23, and 24 gave the
ferrite grain size larger than 6 .mu.m so that their tensile
strengths were below 440 MPa. The steel sheet Nos. 30 and 31 gave
the average carbide grain size larger than 1.2 .mu.m so that their
volume ratio of carbide having larger than 2 .mu.m of the grain
size exceeded 10%, and further their volume ratio of carbide-free
ferrite exceeded 5%, thus the hole-expanding ratio .lambda. was
low, and the stretch-flange formability was poor. The steel sheet
Nos. 21, 25, 28, and 32 gave smaller than 0.1 .mu.m of average
carbide grain size to increase the strength so that the hole
expanding ratio .lambda. and the elongation were low compared with
the target values, and the elongation and the stretch-flange
formability were poor. The steel sheet Nos. 27 and 29 gave larger
than 5% in the volume ratio of carbide-free ferrite so that the
hole expanding ratio .lambda. and the elongation were low compared
with the target values, and the elongation and the stretch-flange
formability were poor. The steel sheet No. 26 gave more than 10% of
the volume ratio of carbide having 2.0 .mu.m or larger grain size,
though the average carbide grain size was in a range from 0.10
.mu.m or larger and smaller than 1.2 .mu.m, thus the hole expanding
ratio .lambda. and the elongation were low compared with the target
values, and the stretch-flange formability and the elongation were
poor.
1TABLE 1 Steel No. C Si Mn P S sol. Al N Other A 0.15 0.22 0.72
0.009 0.005 0.020 0.0038 Cr: 1.0, Mo: 0.16 B 0.23 0.20 0.80 0.010
0.009 0.031 0.0030 -- C 0.35 0.21 0.76 0.014 0.005 0.028 0.0034 --
D 0.35 0.20 0.75 0.012 0.004 0.035 0.0036 Cr: 1.0, Mo: 0.16 E 0.49
0.18 0.75 0.011 0.008 0.030 0.0035 -- F 0.64 0.22 0.73 0.012 0.010
0.021 0.0036 -- G 0.26 0.03 0.45 0.015 0.003 0.040 0.0050 Cr: 0.28
H 0.26 0.03 0.45 0.015 0.003 0.040 0.0050 Mo: 0.30 I 0.47 0.18 0.75
0.011 0.008 0.030 0.0035 Cr: 0.15 J 0.58 0.20 0.74 0.015 0.010
0.021 0.0038 Cr: 0.06 K 0.35 0.21 0.76 0.013 0.005 0.028 0.0034 Cr:
0.18 L 0.35 0.45 0.76 0.013 0.005 0.028 0.0034 Mo: 0.06 M 0.37 0.03
0.75 0.014 0.004 0.028 0.0034 Cr: 0.28, Mo: 0.30 N 0.35 0.18 0.25
0.014 0.005 0.028 0.0034 Mo: 0.15 O 0.35 0.18 0.95 0.014 0.005
0.028 0.0034 Cr: 0.06, Mo: 0.06 P 0.35 0.20 0.75 0.014 0.004 0.031
0.0032 Cr: 0.06, B: 0.0022, Cu: 0.2, Ni: 0.6, W: 0.05 Q 0.34 0.21
0.75 0.013 0.004 0.032 0.0034 Cr: 0.25, Ti: 0.005, Nb: 0.008, V:
0.01, Zr: 0.01 R 0.34 0.21 0.73 0.013 0.004 0.030 0.0038 Cr: 0.06,
Mo: 0.06, Cu: 0.08, Ni: 0.02, Ti: 0.02, V: 0.05
[0106]
2TABLE 2 Primary Primary Rolling cooling Primary cooling Range of
holding Steel termination starting cooling termination temperature
Coiling sheet Steel temperature time rate temperature in the
secondary cooling temperature Annealing No. No. (.degree. C.) (sec)
(.degree. C./sec) (.degree. C.) until the coiling (.degree. C.)
(.degree. C.) condition Remark 1 A Ar3 + 30.degree. C. 0.5 220 590
550.about.590 550 680.degree. C. .times. 40 hr Example 2 B Ar3 +
30.degree. C. 1.2 230 590 570.about.620 580 680.degree. C. .times.
40 hr Example 3 C Ar3 + 20.degree. C. 1.0 210 560 480.about.550 540
680.degree. C. .times. 40 hr Example 4 D Ar3 + 20.degree. C. 1.0
200 550 490.about.530 540 680.degree. C. .times. 40 hr Example 5 E
Ar3 + 30.degree. C. 1.2 200 570 520.about.630 550 710.degree. C.
.times. 40 hr Example 6 F Ar3 + 40.degree. C. 0.4 200 580
580.about.640 560 700.degree. C. .times. 40 hr Example 7 G Ar3 +
20.degree. C. 1.1 210 590 580.about.630 560 680.degree. C. .times.
40 hr Example 8 H Ar3 + 20.degree. C. 1.1 220 580 580.about.620 570
680.degree. C. .times. 40 hr Example 9 I Ar3 + 30.degree. C. 1.2
210 560 530.about.630 560 680.degree. C. .times. 40 hr Example 10 J
Ar3 + 20.degree. C. 1.1 200 570 540.about.620 550 680.degree. C.
.times. 40 hr Example 11 K Ar3 + 20.degree. C. 1.0 210 560
480.about.550 550 680.degree. C. .times. 40 hr Example 12 L Ar3 +
20.degree. C. 1.0 210 570 480.about.570 570 680.degree. C. .times.
40 hr Example 13 M Ar3 + 20.degree. C. 1.0 210 560 480.about.550
560 680.degree. C. .times. 40 hr Example 14 N Ar3 + 20.degree. C.
1.0 210 560 480.about.540 550 680.degree. C. .times. 40 hr Example
15 O Ar3 + 20.degree. C. 1.0 210 570 480.about.550 560 680.degree.
C. .times. 40 hr Example 16 P Ar3 + 20.degree. C. 1.0 210 560
490.about.580 560 680.degree. C. .times. 40 hr Example 17 Q Ar3 +
20.degree. C. 1.0 210 560 500.about.570 560 680.degree. C. .times.
40 hr Example 18 R Ar3 + 20.degree. C. 1.0 210 560 500.about.570
560 680.degree. C. .times. 40 hr Example 19 A Ar3 + 30.degree. C.
0.5 180 680 620.about.650 600 680.degree. C. .times. 40 hr
Comparative Example 20 A Ar3 - 40.degree. C. 1.2 180 590
580.about.630 590 680.degree. C. .times. 40 hr Comparative Example
21 A Ar3 + 10.degree. C. 0.5 280 430 420.about.500 500 660.degree.
C. .times. 40 hr Comparative Example 22 B Ar3 + 30.degree. C. 1.2
210 630 580.about.660 580 680.degree. C. .times. 40 hr Comparative
Example 23 B Ar3 - 40.degree. C. 0.7 160 630 560.about.620 570
700.degree. C. .times. 40 hr Comparative Example 24 B Ar3 +
20.degree. C. 1.2 80 610 550.about.600 540 680.degree. C. .times.
40 hr Comparative Example 25 C Ar3 + 30.degree. C. 0.8 220 580
470.about.550 460 600.degree. C. .times. 40 hr Comparative Example
26 C Ar3 + 20.degree. C. 1.0 210 580 550.about.680 600 680.degree.
C. .times. 40 hr Comparative Example 27 D Ar3 - 30.degree. C. 1.2
160 590 580.about.640 590 680.degree. C. .times. 40 hr Comparative
Example 28 D Ar3 + 20.degree. C. 0.5 280 420 410.about.510 500
660.degree. C. .times. 40 hr Comparative Example 29 E Ar3 -
30.degree. C. 1.2 160 580 550.about.630 520 700.degree. C. .times.
40 hr Comparative Example 30 E Ar3 + 30.degree. C. 0.7 200 660
610.about.650 600 700.degree. C. .times. 40 hr Comparative Example
31 F Ar3 + 20.degree. C. 1.0 180 640 600.about.650 640 700.degree.
C. .times.40 hr Comparative Example 32 F Ar3 + 10.degree. C. 0.6
220 610 540.about.610 560 640.degree. C. .times. 40 hr Comparative
Example
[0107]
3TABLE 3 Average Average Volume ratio of Volume ratio Steel ferrite
carbide carbide larger than of Tensile sheet Steel grain grain 2
.mu.m in grain carbide-free Hardness Hole-expanding strength
Elongation No. No. size (.mu.m) size (.mu.m) size (%) ferrite (HRB)
ratio .lambda. (%) (MPa) (%) Remark 1 A 5.8 0.75 6 5 73 148 440 43
Example 2 B 5.5 0.88 8 5 73 150 445 42 Example 3 C 3.6 0.59 4 3 79
80 490 38 Example 4 D 3.2 0.40 2 3 80 75 500 36 Example 5 E 2.9
0.47 3 2 86 56 560 32 Example 6 F 1.9 0.36 2 1 88 45 590 31 Example
7 G 5.0 0.65 7 4 75 90 470 40 Example 8 H 4.8 0.63 6 4 76 89 480 40
Example 9 I 3.0 0.50 3 2 85 60 550 33 Example 10 J 2.5 0.41 2 1 87
50 580 31 Example 11 K 3.6 0.57 3 3 79 79 490 38 Example 12 L 3.6
0.58 4 4 80 78 500 37 Example 13 M 3.6 0.59 4 3 78 81 480 39
Example 14 N 3.6 0.59 4 3 79 80 490 38 Example 15 O 3.6 0.59 4 3 79
79 490 38 Example 16 P 3.5 0.58 4 3 79 79 490 38 Example 17 Q 3.2
0.58 4 3 80 78 500 37 Example 18 R 3.2 0.59 4 3 79 80 490 38
Example 19 A 10.8 1.44 25 30 70 98 410 42 Comparative Example 20 A
6.8 0.90 9 20 72 118 435 40 Comparative Example 21 A 3.5 0.05 0 1
84 38 535 33 Comparative Example 22 B 6.5 0.94 11 8 72 138 430 40
Comparative Example 23 B 7.2 1.30 15 26 68 75 400 41 Comparative
Example 24 B 6.5 0.88 8 16 72 70 430 40 Comparative Example 25 C
3.4 0.07 0 2 90 21 580 29 Comparative Example 26 C 3.6 1.10 11 5 79
45 490 32 Comparative Example 27 D 5.2 0.64 5 15 78 51 480 33
Comparative Example 28 D 2.1 0.04 0 0 92 20 600 27 Comparative
Example 29 E 3.0 0.68 6 18 82 19 520 28 Comparative Example 30 E
5.2 1.39 22 15 80 20 500 29 Comparative Example 31 F 3.9 1.38 21 6
84 10 530 27 Comparative Example 32 F 3.0 0.08 1 6 89 11 580 25
Comparative Example
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