U.S. patent application number 10/903747 was filed with the patent office on 2005-01-27 for high-ductility steel sheet excellent in press formability and strain age hardenability, and method for manufacturing the same.
This patent application is currently assigned to JFE Steel Corporation, a corporation of Japan. Invention is credited to Furukimi, Osamu, Matsuoka, Saiji, Sakata, Kei, Shimizu, Tetsuo.
Application Number | 20050016644 10/903747 |
Document ID | / |
Family ID | 27346884 |
Filed Date | 2005-01-27 |
United States Patent
Application |
20050016644 |
Kind Code |
A1 |
Matsuoka, Saiji ; et
al. |
January 27, 2005 |
High-ductility steel sheet excellent in press formability and
strain age hardenability, and method for manufacturing the same
Abstract
A steel sheet composition contains appropriate amounts of C, Si,
Mn, P, S, Al and N and 0.5 to 3.0% Cu. A composite structure of the
steel sheet has a ferrite phase or a ferrite phase and a tempered
martensite phase as a primary phase, and a secondary phase
containing retained austenite in a volume ratio of not less than
1%. In place of the Cu, at least one of Mo, Cr, and W may be
contained in a total amount of not more than 2.0%. This composition
is useful in production of a high-ductility hot-rolled steel sheet,
a high-ductility cold-rolled steel sheet and a high-ductility
hot-dip galvanized steel sheet having excellent press formability
and excellent stain age hardenability as represented by a .DELTA.TS
of not less than 80 MPa, in which the tensile strength increases
remarkably through a heat treatment at a relatively low temperature
after press forming.
Inventors: |
Matsuoka, Saiji; (Chiba,
JP) ; Shimizu, Tetsuo; (Kurashiki, JP) ;
Sakata, Kei; (Chiba, JP) ; Furukimi, Osamu;
(Chiba, JP) |
Correspondence
Address: |
IP DEPARTMENT OF PIPER RUDNICK LLP
ONE LIBERTY PLACE, SUITE 4900
1650 MARKET ST
PHILADELPHIA
PA
19103
US
|
Assignee: |
JFE Steel Corporation, a
corporation of Japan
Tokyo
JP
|
Family ID: |
27346884 |
Appl. No.: |
10/903747 |
Filed: |
July 30, 2004 |
Related U.S. Patent Documents
|
|
|
|
|
|
Application
Number |
Filing Date |
Patent Number |
|
|
10903747 |
Jul 30, 2004 |
|
|
|
10163728 |
Jun 6, 2002 |
|
|
|
6818074 |
|
|
|
|
Current U.S.
Class: |
148/651 |
Current CPC
Class: |
C21D 8/0226 20130101;
C23C 2/02 20130101; C21D 8/0278 20130101; C22C 38/16 20130101; C21D
8/0236 20130101; C22C 38/12 20130101; C21D 8/0273 20130101; C22C
38/02 20130101; Y10T 428/12799 20150115; C22C 38/04 20130101; C23C
2/28 20130101 |
Class at
Publication: |
148/651 |
International
Class: |
C21D 008/00 |
Foreign Application Data
Date |
Code |
Application Number |
Jun 6, 2001 |
JP |
JP 2001-170402 |
Jun 29, 2001 |
JP |
JP 2001-198993 |
Jul 3, 2001 |
JP |
JP 2001-202067 |
Claims
1-11. (cancelled)
12. A high-ductility cold-rolled steel sheet excellent in press
formability and in strain age hardenability as represented by a
.DELTA.TS of not less than 80 mpa, comprising a composite structure
containing a primary ferrite phase and a secondary phase containing
a retained austenite phase in a volume ratio of not less than
1%,.
13. A high-ductility steel sheet according to claim 12, wherein the
cold-rolled steel sheet has a composition comprising, in weight
percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not
more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:
not more than 0.3%, N: not more than 0.02%, Cu: 0.5 to 3.0%, and
the balance Fe and incidental impurities.
14. (Original) A high-ductility steel sheet according to claim 13,
the composition further comprising, in weight percent, at least one
of the following Groups A to C, in addition to the above-mentioned
composition: Group A: Ni: not more than 2.0%; Group B: at least one
of Cr and Mo: not more than 2.0% in total; and Group C: at least
one of Nb, Ti, and V: not more than 0.2% in total.
15. A high-ductility steel sheet according to claim 12, wherein the
cold-rolled steel sheet has a composition comprising, in weight
percent: C: not more than 0.20%, Si: not more than 2.0%, Mn: not
more than 3.0% Mn, P: not more than 0.1%, S: not more than 0.02%,
Al: not more than 0.3%, N: not more than 0.02%, at least one
selected from the group consisting of Mo: 05 to 2.0%, Cr: 0.05 to
2.0% and W: 0.05 to 2.0%, not more than 2.0% in total, and the
balance Fe and incidental impurities.
16. A high-ductility steel sheet according to claim 15, the
composition further comprising, in weight percent, at least one of
Nb, Ti, and V, in a total amount of not more than 2.0%.
17. A method for manufacturing a high-ductility cold-rolled steel
sheet excellent in press formability and in strain age
hardenability as typically represented by a .DELTA.TS of not less
than 80 MPa, comprising: a hot rolling step of hot-rolling a steel
slab having a composition containing, in weight percent, C: not
more than 0.20%, Si: not more than 2.0%, Mn: not more than 3.0%, P:
not more than 0.1%, S: not more than 0.02%, Al: not more than 0.3%,
N: not more than 0.02%, and Cu: 0.5 to 3.0% as a material to form a
hot-rolled steel sheet; a cold rolling step of cold-rolling the
hot-rolled steel sheet into a cold-rolled steel sheet; and a
recrystallization annealing step of applying recrystallization
annealing to the cold-rolled steel sheet into a cold-rolled
annealed steel sheet, the recrystallization annealing step
including a heat treatment of heating and soaking the steel sheet
in a ferrite/austenite dual phase region within a temperature range
of the A.sub.C1 transformation point to the A.sub.C3 transformation
point, cooling the sheet, and retaining the sheet in the
temperature region of 300 to 500.degree. C. for 30 to 1,200
seconds.
18. A method for manufacturing a high-ductility cold-rolled steel
sheet according to claim 17, the composition further comprising, in
weight percent, at least one selected from the following Groups A
to C: Group A: Ni: not more than 2.0%; Group B: at least one of Cr
and Mo: not more than 2.0% in total; and Group C: at least one of
Nb, Ti, and V: not more than 0.2% in total.
19. A method for manufacturing a high-ductility cold-rolled steel
sheet according to claim 17, wherein the steel slab is replaced
with a steel slab having a composition containing, in weight
percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not
more than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al:
not more than 0.3%, N: not more than 0.02%, and at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0% and W: 0.05 to 2.0% in a total amount of not more than
2.0%.
20. A method for manufacturing a high-ductility cold-rolled steel
sheet according to claim 19, the composition further comprising, in
weight percent, at least one of Nb, Ti, and V in a total amount of
not more than 2.0%.
21. A method for manufacturing a high-ductility cold-rolled steel
sheet according to claim 17, wherein the hot-rolling step includes
heating the steel slab at a temperature of not less than
900.degree. C., rolling the slab at a finish rolling end
temperature of not less than 700.degree. C., and coiling the
hot-rolled steel sheet at a coiling temperature of not more than
800.degree. C.
22. A method for manufacturing a cold-rolled steel sheet according
to claim 17, wherein all or part of the hot rolling is lubrication
rolling.
23-37. (Cancelled)
Description
BACKGROUND OF THE INVENTION
[0001] 1. Field of the Invention
[0002] The present invention relates mainly to steel sheets for
automobiles, and more particularly, to high-ductility steel sheets
having very high strain age hardenability and excellent press
formability such as ductility, stretch-flanging formability, and
drawability, in which the tensile strength increases remarkably
through a heat treatment after press forming, and to methods for
manufacturing the same. The term "steel sheets" as herein used
shall include hot-rolled steel sheets, cold-rolled steel sheets,
and hot-dip galvanized steel sheets. The term "steel sheets" as
herein used shall also include steel sheets and steel strips.
[0003] 2. Description of the Related Art
[0004] In recent years, weight reduction in automobile bodies has
become a very important issue in relation to emission gas control
for the purpose of preserving global environments. More recently,
efforts are made to achieve higher strength of automotive steel
sheets and to reduce steel sheet thickness in order to reduce the
weights of automobile bodies.
[0005] Because most of the body parts of automobiles made of steel
sheets are formed by press working, steel sheets used must have
excellent press formability. In order to achieve excellent press
formability, it is necessary to ensure high ductility. Stretch
flanging is frequently applied, so that the steel sheets to be used
must have a high hole-expanding ratio. In general, however, a
higher strength of steel sheet tends to result in a lower ductility
and a lower hole-expanding ratio, thus leading to poor press
formability. As a result, there has conventionally been an
increasing demand for high-strength steel sheets having high
ductility and excellent press formability.
[0006] Importance is now placed on safety of an automobile body to
protect a driver and passengers upon collision, and for this
purpose, steel sheets must have improved impact resistance as a
standard of safety upon collision. For the purpose of improving the
crashworthiness, a higher strength in a completed automobile is
more favorable. There has therefore been the strongest demand for
steel sheets having low strength, high ductility, and excellent
press formability upon forming automobile parts, and having high
strength and excellent crashworthiness in completed products.
[0007] To satisfy such a demand, a steel sheet high both in press
formability and strength was developed. This is a bake hardenable
type steel sheet of which the yield stress increases by applying a
bake treatment including holding at a high temperature of 100 to
200.degree. C. after press forming. In this steel sheet, the C
content remaining finally in a solid solution state (solute C
content) is controlled within an appropriate range so as to keep
the softness, shape fixability, and ductility during press forming.
In a bake treatment performed after the press forming of this steel
sheet, the solute C is fixed to a dislocation introduced during the
press forming and inhibits the movement of the dislocation, thus
resulting in an increase in yield stress. In this bake hardenable
type automotive steel sheet, the yield stress can be increased, but
the tensile strength cannot be increased.
[0008] Japanese Examined Patent Application Publication No. 5-24979
discloses a bake hardenable high-strength cold-rolled steel sheet
having a composition comprising C: 0.08 to 0.20%, Mn: 1.5 to 3.5%
and the balance Fe and incidental impurities, and having a
structure composed of uniform bainite containing not more than 5%
of ferrite or composed of bainite partially containing martensite.
The cold-rolled steel sheet disclosed in Japanese Examined Patent
Publication No. 5-24979 is manufactured by rapidly cooling the
steel sheet to a temperature in the range of 400 to 200.degree. C.
in the cooling step after continuous annealing and then slowly
cooling the same. A high degree of baking hardening conventionally
unavailable is thereby achieved through conversion from the
conventional structure mainly comprising ferrite to a structure
mainly comprising bainite in the steel sheet.
[0009] In the steel sheet disclosed in Japanese Examined Patent
Application Publication No. 5-24979, a high degree of baking
hardening conventionally unavailable is obtained through an
increase in yield strength after bake treatment. Even in this steel
sheet, however, it is yet difficult to increase tensile strength
after the bake treatment, and an improvement in crashworthiness
cannot still be achieved.
[0010] On the other hand, some hot-rolled steel sheets are proposed
with a view to increasing not only yield stress but also tensile
strength by applying a heat treatment after press forming.
[0011] For example, Japanese Examined Patent Application
Publication No. 8-23048 proposes a method for manufacturing a
hot-rolled steel sheet comprising the steps of reheating a steel
containing C: 0.02 to 0.13%, Si: not more than 2.0%, Mn: 0.6 to
2.5%, sol. Al: not more than 0.10%, and N: 0.0080 to 0.0250% to a
temperature of not less than 1,100.degree. C. and applying hot
finish rolling at a temperature of 850 to 950.degree. C. The method
also comprising the steps of cooling the hot-rolled steel sheet at
a cooling rate of not less than 15.degree. C./second to a
temperature of less than 150.degree. C., and coiling the same,
thereby forming a composite structure mainly comprising ferrite and
martensite. In the steel sheet manufactured by the technique
disclosed in Japanese Examined Patent Application Publication No.
8-23048, the tensile strength and the yield stress increase by
strain age hardening; however, a serious problem is posed in that
coiling of the steel sheet at a very low coiling temperature as
less than 150.degree. C. results in large variations in mechanical
properties. Another problem includes a large variation in increment
of yield stress after press forming and bake treatments, as well as
poor press formability due to a low hole-expanding ratio (.lambda.)
and decreased stretch-flanging workability.
[0012] Japanese Unexamined Patent Application Publication No.
11-199975 proposes a hot-rolled steel sheet for working excellent
in fatigue characteristics, containing C: 0.03 to 0.20%,
appropriate amounts of Si, Mn, P, S and Al, Cu: 0.2 to 2.0%, and B:
0.0002 to 0.002%, of which the microstructure is a composite
structure comprising ferrite as a primary phase and martensite as a
second phase, and the ferrite phase contains Cu in a solid-solution
and/or precipitation state of not more than 2 nm. The steel sheet
disclosed in Japanese Unexamined Patent Application Publication No.
11-199975 has an object based on the fact that the fatigue limit
ratio is remarkably improved only when Cu and B are added in
combination, and Cu is present in an ultra fine state not more than
2 nm. For this purpose, it is essential to complete hot finish
rolling at a temperature above the A.sub.r3 transformation point,
air-cool the sheet within the temperature region of A.sub.r3 to
A.sub.r1 for 1 to 10 seconds, cool the sheet at a cooling rate of
not less than 20.degree. C./second, and coil the cooled sheet at a
temperature of not more than 350.degree. C. A low coiling
temperature of not more than 350.degree. C. causes serious
deformation of the shape of the hot-rolled steel sheet, thus
inhibiting industrially stable manufacture.
[0013] On the other hand, some automobile parts must have high
corrosion resistance. A hot-dip galvanized steel sheet is suitable
as a material applied to portions requiring high corrosion
resistance. For this reason, a particular demand exists for hot-dip
galvanized steel sheets excellent in press formability during
forming, and is considerably hardened by a heat treatment after the
forming.
[0014] To respond to such a demand, for example, Japanese Patent
Publication No. 2802513 proposes a method for manufacturing a
hot-dip galvanized steel sheet using a hot-rolled steel sheet as a
black plate. The method comprises the steps of hot-rolling a steel
slab containing C: not more than 0.05%, Mn: 0.05 to 0.5%, Al: not
more than 0.1% and Cu: 0.8 to 2.0% at a coiling temperature of not
more than 530.degree. C. The method further comprising the
subsequent steps of reducing the steel sheet surface by heating the
hot-rolled steel sheet to a temperature of not more than
530.degree. C., and hot-dip-galvanizing the sheet, whereby
remarkable hardening is available through a heat treatment after
forming. In the steel sheet manufactured by this method, however,
the heat treatment temperature must be high as not less than
500.degree. C., in order to obtain remarkable hardening from the
heat treatment after the forming, and this has a problem in
practice.
[0015] Japanese Unexamined Patent Application Publication No.
10-310824 proposes a method for manufacturing an alloyed hot-dip
galvanized steel sheet having increased strength by a heat
treatment after forming, using a hot-rolled or cold-rolled steel
sheet as a black plate. This method comprises the steps of
hot-rolling a steel containing C: 0.01 to 0.08%, appropriate
amounts of Si, Mn, P, S, Al and N, and at least one of Cr, W and
Mo: 0.05 to 3.0% in total. The method further comprises the step of
cold-rolling or temper-rolling and annealing the sheet. The method
still further comprises the step of applying hot-dip galvanizing to
the sheet and heating the sheet for alloying treatment. The tensile
strength of the steel sheet is increased by heating the sheet at a
temperature within the range of 200 to 450.degree. C. However, the
resultant steel sheet involves a problem in that the microstructure
comprises a ferrite single phase, a ferrite and pearlite composite
structure, or a ferrite and bainite composite structure; hence,
high ductility and low yield strength are unavailable, resulting in
low press formability.
SUMMARY OF THE INVENTION
[0016] The present invention was made in view of the fact that, in
spite of the strong demand as described above, a technique for
industrially stably manufacturing a steel sheet satisfying these
properties has never been found. The present invention solves the
problems described above. It is an object of the present invention
to provide is directed to high-ductility and high-strength steel
sheets suitable for automobiles and having excellent press
formability and excellent strain age hardenability, in which the
tensile strength increases considerably through a heat treatment at
a relatively low temperature after press forming. It is also an
object of the present invention to provide a manufacturing method
capable of stably manufacturing the high-ductility and
high-strength steel sheets.
[0017] To achieve the above-mentioned object of the invention, the
inventors carried out extensive studies on the effect of the steel
sheet structure and alloying elements on strain age hardenability.
As a result, the inventors found that a steel sheet having high age
hardenability which leads to both an increase in yield stress and a
remarkable increase in tensile strength can be obtained after a
pre-deformation treatment with a prestrain of not less than 5% and
a heat treatment at a relatively low temperature as within the
range of 150 to 350.degree. C. by (1) forming a composite structure
of the steel sheet comprising ferrite and a phase containing
retained austenite in a volume ratio of not less than 1%, and (2)
limiting the C content within the range of a low-carbon region to a
medium-carbon region and containing Cu within an appropriate range
or at least one of Mo, Cr, and W in place of Cu. In addition, the
steel sheet was found to have satisfactory ductility, a high hole
expanding ratio, and excellent press formability.
[0018] The results of a fundamental experiment carried out by the
inventors on hot-rolled steel sheets will first be described.
[0019] A sheet bar having a composition comprising, in weight
percent, C: 0.10%, Si: 1.4%, Mn: 1.5%, P: 0.01%, S: 0.005%, Al:
0.04%, N: 0.002% and Cu: 0.3 or 1.3% was heated to 1,250.degree. C.
and soaked. Then, the sheet bar was subjected to three-pass rolling
into a thickness of 2.0 mm so that the finish rolling end
temperature was 850.degree. C. Thereafter, cooling conditions and
the coiling temperature were changed variously to convert a single
ferrite structure steel sheet into a hot-rolled steel sheet with a
composite structure composed of ferrite as a primary phase and a
retained austenite-containing phase as a secondary phase
(hereinafter, referred to also as a composite ferrite/retained
austenite structure).
[0020] Tensile properties were investigated by a tensile test on
the resultant hot-rolled steel sheets. A pre-deformation treatment
of a tensile prestrain of 5% was applied to each test piece sampled
from these hot-rolled steel sheets. Then, after applying a heat
treatment at 50 to 350.degree. C. for 20 minutes, a tensile test
was carried out to determine tensile properties, and the strain age
hardenability was evaluated.
[0021] The strain age hardenability was evaluated in terms of the
increment .DELTA.TS that is a difference between the tensile
strength TS.sub.HT after heat treatment and the tensile strength TS
before the heat treatment. That is, .DELTA.TS=(tensile strength
TS.sub.HT after heat treatment)-(tensile strength TS before
pre-deformation treatment). The tensile test was carried out by
using JIS No. 5 tensile test pieces sampled in the rolling
direction.
[0022] FIG. 1 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the steel sheet structure. A
pre-deformation treatment of a tensile prestrain of 5% and then a
heat treatment of 250.degree. C. for 20 minutes were applied to the
test pieces. The increment .DELTA.TS was determined from the
difference in tensile strength TS between before and after the heat
treatment. FIG. 1 suggests that, for a Cu content of 1.3 wt. %, a
high strain age hardenability as represented by a .DELTA.TS of not
less than 80 MPa is obtained by forming a composite
ferrite/retained austenite steel sheet structure. For a Cu content
of 0.3 wt. %, .DELTA.TS is less, than 80 MPa, irrespective of the
steel sheet structure, and high strain age hardenability cannot be
obtained.
[0023] It is possible to manufacture a hot-rolled steel sheet
having a high strain age hardenability by limiting the Cu content
within an appropriate range, and forming a composite structure
having ferrite as a primary phase and a retained
austenite-containing phase as a secondary phase.
[0024] FIG. 2 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the heat treatment temperature
after pre-strain treatment. The microstructure of the steel sheet
is a composite structure having ferrite as a primary phase and a
retained austenite-containing phase as a secondary phase, and the
volume ratio of the retained austenite structure is 8% of the
entire structure.
[0025] FIG. 2 shows that the increment .DELTA.TS increases as the
heat treatment temperature increases and strongly depends on the Cu
content. With a Cu content of 1.3 wt. %, a high strain age
hardenability as represented by a .DELTA.TS of not less than 80 MPa
is obtained at a heat treatment temperature of not less than
150.degree. C. For a Cu content of 0.3 wt. %, .DELTA.TS is less
than 80 MPa at any heat treatment temperature, and high strain age
hardenability cannot be obtained.
[0026] In addition, a hole expanding test was carried out on steel
sheets having a single ferrite structure or a composite
ferrite/retained austenite structure, and Cu contents of 0.3 wt %
and 1.3 wt %, and the hole expanding ratio .lambda. was determined.
In the hole expanding test, punch holes were formed in test pieces
through punching with a punch having a diameter of 10 mm.
Thereafter, hole expansion was conducted with a conical punch
having a vertical angle of 60 degrees so that the burr was outside,
until cracks passing through the sheet in the thickness direction
form. The hole expanding ratio .lambda. was determined by the
formula: .lambda.(%)={(d-d.sub.0)/d.sub.0}.times.100 where d.sub.0
represents the initial hole diameter, and d represents the hole
inside diameter on occurrence of cracks.
[0027] In the case of a Cu content of 1.3 wt %, a hot-rolled steel
sheet having a composite ferrite/retained austenite structure had a
hole expanding ratio of about 140%, and a hot-rolled steel sheet
having a single ferrite structure also had a hole expanding ratio
of about 140%. In contrast, in the case of a Cu content of 0.3 wt
%, a hot-rolled steel sheet having a single ferrite structure had a
hole expanding ratio of 120%, and a hot-rolled steel sheet having a
composite ferrite/retained austenite structure had a hole expanding
ratio of about 80%.
[0028] As described above, it is clear that the hot-rolled steel
sheet having a composite ferrite/retained austenite structure has
an increased hole expanding ratio and that hole expanding
formability is improved with an increased Cu content. A detailed
mechanism of the improvement in hole expanding formability by Cu
has not yet been clarified. The contained Cu is considered to
reduce the difference in hardness between the ferrite/retained
austenite and the strain-induced transformed martensite.
[0029] In the hot-rolled steel sheet of the present invention, very
fine Cu precipitates in the steel sheet as a result of a
pre-deformation with a strain of 2% or more as measured upon
measuring the increment of deformation stress from before to after
a usual heat treatment and the heat treatment carried out at a
relatively low temperature in the range of 150 to 350.degree. C.
According to a study carried out by the present inventors, high
strain age hardenability bringing about an increase in yield stress
and a remarkable increase in tensile strength probably achieved by
the precipitation of very fine Cu. Such precipitation of very fine
Cu by a heat treatment in a low-temperature region has never been
observed in ultra-low carbon steel or low-carbon steel in reports
so far released. A reason for precipitation of very fine Cu in a
heat treatment at a low temperature has not as yet been clarified
to date. However, it is presumable as follows. During isothermal
holding in the temperature range of 620 to 780.degree. C. or during
slow cooling from this temperature range after rapid cooling
subsequent to hot rolling, a large amount of Cu is distributed to
the y phase. After cooling, Cu is dissolved in the retained
austenite in a supersaturation state. The retained austenite is
transformed into martensite by a prestrain of not less than 5%, and
very fine Cu precipitates in the strain-induced transformed
martensite during a subsequent low-temperature treatment.
[0030] Next, the results of a fundamental experiment carried out by
the present inventors on the cold-rolled steel sheet will be
described.
[0031] A sheet bar having a composition comprising, in weight
percent, C: 0.10%, Si: 1.2%, Mn: 1.4%, P: 0.01%, S: 0.005%, Al:
0.03%, N: 0.002%, and Cu: 0.3 or 1.3% was heated to 1,250.degree.
C., soaked and subjected to three-pass rolling into a thickness of
4.0 mm so that the finish rolling end temperature was 900.degree.
C. After the completion of finish rolling, a temperature holding
equivalent treatment of 600.degree. C. for 1 hour was applied as a
coiling treatment. Thereafter, the sheet was cold-rolled at a
reduction of 70% into a cold-rolled steel sheet having a thickness
of 1.2 mm. Then, the cold-rolled sheet was heated at a temperature
in the range of 700 to 850.degree. C. and soaked for 60 seconds.
Thereafter, the sheet was cooled to 400.degree. C., and was held at
the temperature (400.degree. C.) for 300 seconds for
recrystallization annealing. By the recrystallization annealing,
various cold-rolled steel sheets were obtained in which the
structure changed from a single ferrite structure to a composite
ferrite/retained austenite structure.
[0032] Tensile tests were conducted on the resultant cold-roll
steel sheets as in the hot-rolled steel sheets to determine tensile
properties. Tensile properties (YS, TS) were determined by sampling
test pieces from these cold-rolled steel sheets, applying a
pre-deformation treatment with a tensile prestrain of 5% to these
test pieces, then heating the steel sheets at 50 to 350.degree. C.
for 20 minutes, and then conducting the tensile tests.
[0033] The strain age hardenability was evaluated in terms of the
tensile strength increment .DELTA.TS from before to after the heat
treatment, as in the hot-rolled steel sheet.
[0034] FIG. 3 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the recrystallization annealing
temperature. The value .DELTA.TS was determined by applying a
pre-deformation treatment with a tensile prestrain of 5% to test
pieces sampled from the resultant cold-rolled steel sheets,
conducting a heat treatment of 250.degree. C. for 20 minutes, and
carrying out a tensile test.
[0035] FIG. 3 suggests that a high strain age hardenability as
represented by a .DELTA.TS of not less than 80 MPa is available, in
the case of a Cu content of 1.3 wt. %, by employing a
recrystallization annealing temperature of not less than
750.degree. C. to convert the steel sheet structure into a
composite ferrite/retained austenite structure. On the other hand,
in the case of a Cu content of 0.3 wt. %, high strain age
hardenability is unavailable because .DELTA.TS is less than 80 MPa
at any recrystallization annealing temperature. FIG. 3 suggests the
possibility of manufacturing a cold-rolled steel sheet having a
high strain age hardenability by optimizing the Cu content and
forming a composite ferrite/retained austenite structure.
[0036] FIG. 4 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the heat treatment temperature
after pre-strain treatment. The steel sheet used was annealed at
800.degree. C., which was the dual phase region of ferrite
(.alpha.)+austenite (.gamma.), for a holding time of 60 seconds
after cold rolling, cooled from the holding temperature
(800.degree. C.) to 400.degree. C. at a cooling rate of 30.degree.
C./second, and held at 400.degree. C. for 300 seconds. The steel
sheets had a composite ferrite/retained austenite (secondary phase)
microstructure, the volume ratio of the retained austenite
structure being 4%.
[0037] FIG. 4 shows that the increment .DELTA.TS increases as the
heat treatment temperature increases and strongly depends on the Cu
content. With a Cu content of 1.3 wt. %, a high strain age
hardenability as represented by a .DELTA.TS of not less than 80 MPa
is obtained at a heat treatment temperature of not less than
150.degree. C. For a Cu content of 0.3 wt. %, .DELTA.TS is less
than 80 MPa at any heat treatment temperature, and high strain age
hardenability cannot be obtained.
[0038] In addition, a hole expanding test was carried on
cold-rolled steel sheets having a composite ferrite/retained
austenite structure and Cu contents of 0.3 wt % and 1.3 wt. % to
determine the hole expanding ratio (.lambda.), as in the hot-rolled
steel sheet.
[0039] In the cold-rolled steel sheet with a Cu content of 1.3%,
.lambda. was 130%; while in the cold-rolled steel sheet with a Cu
content of 0.3%, .lambda. was 60%. It is clear that, for a Cu
content of 1.3 wt. %, the hole expanding ratio is increased and
hole expanding formability is improved even in the cold-rolled
steel sheet, as in the hot-rolled steel sheet. A detailed mechanism
of improvement in hole expanding formability with content of Cu has
not yet been clarified, as in the hot-rolled steel sheet. Also, in
the cold-rolled steel sheet, it is considered that the contained Cu
reduces the difference in hardness between the ferrite/retained
austenite structure and the strain-induced transformed martensite
structure.
[0040] In the cold-rolled steel sheet of the present invention,
very fine Cu precipitates in the steel sheet as a result of a
pre-deformation with a strain larger than 2%, which is equivalent
to the prestrain on measuring the deformation stress increment from
before to after a usual heat treatment, and a heat treatment at a
relatively low temperature of 150 to 350.degree. C. According to a
study carried out by the present inventors, also in the cold-rolled
steel sheet, high strain age hardenability bringing about an
increase in yield stress and a remarkable increase in tensile
strength is probably achieved by the precipitation of very fine Cu.
A reason for precipitation of very fine Cu in a heat treatment in a
low temperature region has not as yet been clarified to date.
However, it is presumable as follows. During recrystallization
annealing in the dual phase region of .alpha.+.gamma., a large
amount of Cu is distributed to the y phase. The distributed Cu
remains even after cooling and is dissolved into the martensite in
a supersaturation state, and very fine Cu precipitates through a
prestrain of not less than 5% and a low-temperature treatment.
[0041] Next, the result of a fundamental experiment carried out by
the present inventors on the hot-dip galvanized steel sheet will be
described.
[0042] A sheet bar having a composition comprising, in weight
percent, C: 0.08%, Si: 0.5%, Mn: 2.0%, P: 0.01%, S: 0.004%, Al:
0.04%, N: 0.002% and Cu: 0.3 or 1.3% was heated to 1,250.degree. C.
and soaked. Then, the sheet bar was subjected to three-pass rolling
into a thickness of 4.0 mm so that the finish rolling end
temperature was 900.degree. C. After the finish rolling, a
temperature holding equivalent treatment of 600.degree. C. for 1 h
was applied as a coiling treatment. Thereafter, the hot-rolled
sheet was cold-rolled at a reduction of 70% into a cold-rolled
steel sheet having a thickness of 1.2 mm. Then, the cold-rolled
sheet was heated and soaked at 900.degree. C., and cooled at a
cooling rate of 30.degree. C./sec. (a primary heat treatment) . The
steel sheet after the primary heat treatment had a lath martensite
structure. The steel sheet after the primary heat treatment was
subjected to a secondary heat treatment at various temperatures,
then rapidly cooled to a temperature in the range of 450 to
500.degree. C. Then, the sheet was immersed into a hot-dip
galvanizing bath (0.13 wt. % Al--Zn bath) to form a hot-dip
galvanizing layer on the surface. Further, the sheet was reheated
to a temperature in the range of 450 to 550.degree. C. to alloy the
hot-dip galvanizing layer (Fe content in the galvanizing layer:
about 10%).
[0043] For the resultant hot-dip galvanized steel sheet, tensile
properties were determined through a tensile test. In addition,
test pieces were sampled from the hot-dip galvanized steel sheet,
and a pre-deformation treatment with a tensile prestrain of 5% was
applied to the test pieces, as in the hot-rolled steel sheet and
the cold-rolled steel sheet. Then, a heat treatment of 50 to
350.degree. C. for 20 minutes was applied. Thereafter, a tensile
test was carried out to determine tensile properties. The strain
age hardenability was evaluated in terms of the increment .DELTA.TS
of the tensile strength from before to after the heat
treatment.
[0044] FIG. 5 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the secondary heat treatment
temperature. The increment .DELTA.TS was determined by applying a
tensile prestrain of 5% to test pieces sampled from the resultant
hot-dip galvanized steel sheets, conducting a heat treatment at
250.degree. C. for 20 minutes, and carrying out a tensile test.
[0045] FIG. 5 suggests that, for a Cu content of 1.3 wt. %, a high
strain age hardenability as represented by a .DELTA.TS of not less
than 80 MPa can be obtained by forming a composite ferrite/tempered
martensite/retained austenite steel sheet structure. In contrast,
in the case of a Cu content of 0.3 wt. %, high strain age
hardenability cannot be obtained as because .DELTA.TS is less than
80 MPa at any secondary heat treatment temperature.
[0046] FIG. 5 suggests the possibility of manufacturing a hot-dip
galvanized steel sheet having high strain age hardenability by
optimizing the Cu content and by forming a composite
ferrite/tempered martensite/retained austenite structure.
[0047] FIG. 6 illustrates the effect of the Cu content on the
relationship between .DELTA.TS and the heat treatment temperature
after pre-strain treatment. The increment .DELTA.TS was determined
by applying a tensile prestrain of 5% to test pieces sampled from
the alloyed hot-dip galvanized steel sheets treated at a secondary
heat treatment temperature of 800.degree. C., conducting a heat
treatment of 50 to 350 .degree. C. for 20 minutes, and carrying out
a tensile test.
[0048] FIG. 6 shows that the increment .DELTA.TS increases as the
heat treatment temperature increases after the pre-deformation
treatment and strongly depends on the Cu content. With a Cu content
of 1.3 wt. %, a high strain age hardenability as represented by a
.DELTA.TS of not less than 80 MPa can be obtained at a heat
treatment temperature of not less than 150.degree. C. In contrast,
for a Cu content of 0.3 wt. %, .DELTA.TS is less than 80 MPa at any
heat treatment temperature, and high strain age hardenability
cannot be obtained.
[0049] In the hot-dip galvanized steel sheet of the present
invention, very fine Cu precipitates in the steel sheet as a result
of a pre-deformation with a strain larger than 2% which is a usual
amount of strain on measuring the deformation stress increment from
before to after a heat treatment, and a heat treatment within a
relatively low temperature region of 150 to 350.degree. C.
According to a study carried out by the present inventors, high
strain age hardenability bringing about an increase in yield stress
and a remarkable increase in tensile strength is probably achieved
by the precipitation of very fine Cu. A reason for precipitation of
very fine Cu in a heat treatment in a low temperature region has
not as yet been clarified to date. However, it is presumable as
follows. During heat treatment in the dual phase region of ferrite
(.alpha.)+austenite (.gamma.), a large amount of Cu is distributed
to the .gamma. phase, and the distributed Cu remaining even after
cooling is dissolved into the retained austenite in a
supersaturation state. The retained austenite is transformed into
martensite by a prestrain of not less than 5%, and very fine Cu
precipitates in the martensite through a subsequent low-temperature
heat treatment.
[0050] In addition, hole expanding test was performed using.
hot-dip galvanized steel sheets having a composite structure of
ferrite/tempered martensite/retained austenite and Cu contents of
0.3 wt % and 1.3 wt. % to determine the hole expanding ratio
(.lambda.), as in the hot-rolled steel sheet and the cold-rolled
steel sheet.
[0051] The hole expanding ratio .lambda. of the steel sheet having
a Cu content of 1.3% was 120%, while the hole expanding ratio
.lambda. of the steel sheet having a Cu content of 0.3% was 50%.
The results suggest that for a Cu content of 1.3 wt %, the hole
expanding ratio is increased and hole expanding formability is
improved, as compared with a Cu content of 0.3%.
[0052] A detailed mechanism of improvement in hole expanding
formability with content of Cu has not yet been clarified, as in
the hot-rolled steel sheet and the cold-rolled steel sheet, but it
is considered that the contained Cu reduces the difference in
hardness among the ferrite, the tempered martensite/retained
austenite, and the martensite formed by strain induced
transformation.
[0053] On the basis of the novel findings as described above, the
present inventors carried out further extensive studies and found
that the above-mentioned phenomena occurred in steel sheets not
containing Cu as well.
[0054] The structure of a steel sheet having a composition
containing at least one of Mo, Cr, and W was converted to a
composite structure containing a ferrite primary phase and a phase
containing retained austenite as a secondary phase. Thereafter, by
applying a prestrain and a heat treatment in a low temperature
region, it was found that very fine carbides precipitated in the
strain-induced transformed martensite, resulting in an increase in
tensile strength. The strain-induced fine precipitation at a low
temperature was more remarkable in a steel composition containing
at least one of Nb, Ti, and V in addition to at least one of Mo,
Cr, and W.
[0055] The present invention was completed through further studies
on the basis of the aforementioned findings. The gist of the
present invention is as follows:
[0056] (1) A high-ductility steel sheet excellent in press
formability and in strain age hardenability as represented by a
.DELTA.TS of not less than 80 MPa, comprising a composite structure
containing a primary phase containing a ferrite phase and a
secondary phase containing a retained austenite phase in a volume
ratio of not less than 1%.
[0057] (2) A high-ductility steel sheet according to aspect (1),
wherein the steel sheet is a hot-rolled steel sheet, and the
primary phase consisting essentially of a ferrite phase.
[0058] (3) A high-ductility steel sheet according to aspect (2),
wherein the hot-rolled steel sheet has a composition comprising, in
weight percent, C: 0.05 to 0.20%, Si: 1.0 to 3.0%, Mn: not more
than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al: not
more than 0.30%, N: not more than 0.02%, and Cu: 0.5 to 3.0%, and
the balance Fe and incidental impurities.
[0059] (4) A high-ductility steel sheet according to aspect (3),
the composition further comprising, in weight percent, at least one
of the following Groups A to C:
[0060] Group A: Ni: not more than 2.0%;
[0061] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0062] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0063] (5) A high-ductility steel sheet according to aspect (2),
wherein the hot-rolled steel sheet has a composition comprising, in
weight percent, C: 0.05 to 0.20%, Si: 1.0 to 3.0%, Mn: not more
than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al: not
more than 0.30%, N: not more than 0.02%, at least one of Mo: 0.05
to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0%, not more than 2.0%
in total, and the balance Fe and incidental impurities.
[0064] (6) A high-ductility steel sheet according to aspect (5),
the composition further containing, in weight percent, at least one
of Nb, Ti, and V in an amount of not more than 2.0% in total.
[0065] (7) A method for manufacturing a high-ductility hot-rolled
steel sheet excellent in press formability and in strain age
hardenability as represented by a .DELTA.TS of not less than 80
MPa, comprising the steps of: hot-rolling a steel slab having a
composition comprising, in weight percent, C: not more than 0.20%,
Si: 1.0 to 3.0%, Mn: not more than 3.0%, P: not more than 0.10%, S:
not more than 0.02%, Al: not more than 0.30%, N: not more than
0.02%, and Cu: 0.5 to 3.0%, into a hot-rolled steel sheet having a
prescribed thickness, the hot rolling step including finish-rolling
at a finish rolling end temperature of 780 to 980.degree. C.;
cooling the finish-rolled steel sheet to a temperature in the range
of 620 to 780.degree. C. within 2 seconds at a cooling rate of at
least 50.degree. C./second; holding the sheet at the temperature in
the range of 620 to 780.degree. C. for 1 to 10 seconds, or slowly
cooling the sheet at a cooling rate of not more than 20.degree.
C./second; cooling the sheet at a cooling rate of not less than
50.degree. C./second to a temperature of 300 to 500.degree. C.; and
coiling the sheet.
[0066] (8) A method for manufacturing a high-ductility hot-rolled
steel sheet excellent in press formability and in strain age
hardenability as represented by a .DELTA.TS of at least 80 MPa,
according to aspect (7), the composition further comprising, in
weight percent, at least one of the following Groups A to C:
[0067] Group A: Ni: not more than 2.0%;
[0068] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0069] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0070] (9) A method for manufacturing a high-ductility hot-rolled
steel sheet according to aspect (7), wherein the steel slab is
replaced with a steel slab having a composition containing, in
weight percent, C: 0.05 to 0.20%, Si: 1.0 to 3.0%, Mn: not more
than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al: not
more than 0.30%, N: not more than 0.02%, and at least one of Mo:
0.05 to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0% in a total
amount of not more than 2.0%.
[0071] (10) A method for manufacturing a high-ductility hot-rolled
steel sheet according to aspect (9), the composition further
containing, in weight percent, at least one of Nb, Ti, and V in a
total amount of not more than 2.0%.
[0072] (11) A method for manufacturing a high-ductility hot-rolled
steel sheet according to any one of aspects (7) to (10), wherein
all or part of the finish rolling is lubrication rolling.
[0073] (12) A high-ductility steel sheet according to aspect (1),
wherein the steel sheet is a cold-rolled steel sheet, and the
primary phase containing the ferrite phase is a ferrite phase.
[0074] (13) A high-ductility steel sheet according to aspect (12),
wherein the cold-rolled steel sheet has a composition comprising,
in weight percent, C: not more than 0.20%, Si: not more than 2.0%,
Mn: not more than 3.0%, P: not more than 0.1%, S: not more than
0.02%, Al: not more than 0.3%, N: not more than 0.02%, Cu: 0.5 to
3.0%, and the balance Fe and incidental impurities.
[0075] (14) A high-ductility steel sheet according to aspect (13),
the composition further comprising, in weight percent, at least one
of the following Groups A to C:
[0076] Group A: Ni: not more than 2.0%;
[0077] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0078] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0079] (15) A high-ductility steel sheet according to aspect (12),
wherein the cold-rolled steel sheet has a composition comprising,
in weight percent: C: not more than 0.20%, Si: not more than 2.0%,
Mn: not more than 3.0%, P: not more than 0.1%, S: not more than
0.02%, Al: not more than 0.3%, N: not more than 0.02%, at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0% and W: 0.05 to 2.0%, not more than 2.0% in total, and the
balance Fe and incidental impurities.
[0080] (16) A high-ductility steel sheet according to aspect (15),
the composition further comprising, in weight percent, at least one
of Nb, Ti, and V, in a total amount of not more than 2.0%.
[0081] (17) A method for manufacturing a high-ductility cold-rolled
steel sheet excellent in press formability and in strain age
hardenability as represented by a .DELTA.TS of not less than 80
MPa, comprising: a hot rolling step of hot-rolling a steel slab
having a composition containing, in weight percent, C: not more
than 0.20%, Si: not more than 2.0%, Mn: not more than 3.0%, P: not
more than 0.1%, S: not more than 0.02%, Al: not more than 0.3%, N:
not more than 0.02%, and Cu: 0.5 to 3.0% as a material to form a
hot-rolled steel sheet; a cold rolling step of cold-rolling the
hot-rolled steel sheet into a cold-rolled steel sheet; and a
recrystallization annealing step of applying recrystallization
annealing to the cold-rolled steel sheet into a cold-rolled
annealed steel sheet, the recrystallization annealing step
including a heat treatment of heating and soaking the steel sheet
in a ferrite/austenite dual phase region within a temperature range
of the A.sub.C1 transformation point to the A.sub.C3 transformation
point, cooling the sheet, and holding the sheet in the temperature
region of 300 to 500.degree. C. for 30 to 1,200 seconds.
[0082] (18) A method for manufacturing a high-ductility cold-rolled
steel sheet according to aspect (17), the composition further
containing, in weight percent, at least one selected from the
following Groups A to C:
[0083] Group A: Ni: not more than 2.0%;
[0084] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0085] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0086] (19) A method for manufacturing a high-ductility cold-rolled
steel sheet according to aspect (17), wherein the steel slab is
replaced with a steel slab having a composition containing, in
weight percent, C: not more than 0.20%, Si: not more than 2.0%, Mn:
not more than 3.0%, P: not more than 0.10%, S: not more than 0.02%,
Al: not more than 0.3%, N: not more than 0.02%, and at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0% and W: 0.05 to 2.0% in a total amount of not more than
2.0%.
[0087] (20) A method of manufacturing a high-ductility cold-rolled
steel sheet according to aspect (19), the composition further
containing, in weight percent, at least one of Nb, Ti, and V in a
total amount of not more than 2.0%.
[0088] (21) A method for manufacturing a high-ductility cold-rolled
steel sheet according to any one of aspects (17) to (20), wherein
the hot-rolling step includes heating the steel slab at a
temperature of not less than 900.degree. C., rolling the slab at a
finish rolling end temperature of not less than 700.degree. C., and
coiling the hot-rolled steel sheet at a coiling temperature of not
more than 800.degree. C.
[0089] (22) A method for manufacturing a cold-rolled steel sheet
according to any one of aspects (17) to (21), wherein all or part
of the hot rolling is lubrication rolling.
[0090] (23) A high-ductility hot-dip galvanized steel sheet
comprising a hot-dip galvanizing layer or an alloyed hot-dip
galvanizing layer formed on the surface of the high-ductility steel
sheet according to any one of aspects (1) to (6).
[0091] (24) A high-ductility hot-dip galvanized steel sheet
comprising a hot-dip galvanizing layer or an alloyed hot-dip
galvanizing layer formed on the surface of the high-ductility steel
sheet according to any one of aspects (12) to (16).
[0092] (25) A high-ductility steel sheet according to aspect (1)
wherein the steel sheet is a hot-dip galvanized steel sheet having
a hot-dip galvanizing layer or an alloyed hot-dip galvanizing layer
formed on a surface of the steel sheet, and the primary phase
containing a ferrite phase comprises a ferrite phase and a tempered
martensite phase.
[0093] (26) A high-ductility steel sheet according to aspect (25),
wherein the steel sheet has a composition comprising, in weight
percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not
more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:
not more than 0.3%, N: not more than 0.02%, Cu: 0.5 to 3.0%, and
the balance Fe and incidental impurities.
[0094] (27) A high-ductility steel sheet according to aspect (26),
the composition further containing, in weight percent, at least one
of the following Groups A to C:
[0095] Group A: Ni: not more than 2.0%;
[0096] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0097] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0098] (28) A high-ductility steel sheet according to aspect (25),
wherein the steel sheet has a composition comprising, in weight
percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not
more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:
not more than 0.3%, N: not more than 0.02%, at least one selected
from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0% and
W: 0.05 to 2.0% in a total amount of not more than 2.0%, and the
balance Fe and incidental impurities.
[0099] (29) A high-ductility steel sheet according to aspect (28),
the composition further containing, in weight percent, at least one
of Nb, Ti, and V in a total amount of not more than 2.0%.
[0100] (30) A method for manufacturing of a high-ductility hot-dip
galvanized steel sheet excellent in press formability and in strain
age hardenability as represented by a .DELTA.TS of not less than 80
MPa, comprising: a primary heat-treating step of heating a steel
sheet to a temperature of not less than the A.sub.C1 transformation
point and rapidly cooling the steel sheet, the steel sheet having a
composition containing, in weight percent, C: not more than 0.20%,
Si: not more than 2.0%, Mn:
[0101] not more than 3.0%, P: not more than 0.1%, S: not more than
0.02%, Al: not more than 0.3%, N: not more than 0.02%, and Cu: 0.5
to 3.0%; a secondary heat-treating step of heating the steel sheet
to a temperature in the range of the A.sub.C1 transformation point
to the A.sub.C3 transformation point; and a hot-dip galvanizing
step of forming a hot-dip galvanizing layer on the surface of the
steel sheet.
[0102] (31) A method for manufacturing a high-ductility cold-rolled
steel sheet according to aspect (30), the composition further
containing, in weight percent, at least one of the following Groups
A to C:
[0103] Group A: Ni: not more than 2.0%;
[0104] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0105] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0106] (32) A method for manufacturing a high-ductility hot-dip
galvanized steel according to aspect (30), wherein the steel sheet
is replaced with a steel sheet having a composition comprising, in
weight percent, C: not more than 0.20%, Si: not more than 2.0%, Mn:
not more than 3.0%, P: not more than 0.1%, S: not more than 0.02%,
Al: not more than 0.3%, N: not more than 0.02%, and at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0% and W: 0.05 to 2.0% in a total amount of not more than
2.0%.
[0107] (33) A method for manufacturing a high-ductility hot-dip
galvanized steel sheet according to aspect (32), the composition
further containing, in weight percent, at least one of Nb, Ti, and
V in a total amount of not more than 2.0%.
[0108] (34) A method for manufacturing a high-ductility hot-dip
galvanized steel sheet according to any one of aspects (30) to
(33), further comprising a pickling treatment step of pickling the
steel sheet between the primary heat-treating step and the
secondary heat-treating step.
[0109] (35) A method for manufacturing a high-ductility hot-dip
galvanized steel sheet according to any one of aspects (30) to
(34), further comprising an alloying step of alloying the hot-dip
galvanizing layer, subsequent to the hot-dip galvanizing step.
[0110] (36) A method for manufacturing a high-strength hot-dip
galvanized steel sheet according to any one of aspects (30) to
(35), wherein the steel sheet is a hot rolled steel sheet
manufactured by hot-rolling a material under conditions including a
heating temperature of not less than 900.degree. C., a finish
rolling end temperature of not less than 700.degree. C. and a
coiling temperature of not more than 800.degree. C., or a
cold-rolled steel sheet obtained by cold-rolling the hot-rolled
steel sheet.
[0111] (37) A method for manufacturing a high-strength hot-dip
galvanized steel sheet according to aspect (36), wherein the
cold-rolling is performed at a reduction ratio of not less than
40%.
BRIEF DESCRIPTION OF THE DRAWINGS
[0112] FIG. 1 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the steel sheet structure
after a pre-deformation and a heat treatment of a hot-rolled steel
sheet;
[0113] FIG. 2 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the heat treatment
temperature after a pre-deformation and a heat treatment of a
hot-rolled steel sheet;
[0114] FIG. 3 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the recrystallization
annealing temperature after pre-deformation and a heat treatment of
a cold-rolled steel sheet;
[0115] FIG. 4 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the heat treatment
temperature after pre-deformation and a heat treatment of a
cold-rolled steel sheet;
[0116] FIG. 5 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the secondary heat
treatment temperature after a pre-deformation and a heat treatment
of a hot-dip galvanized steel sheet; and
[0117] FIG. 6 is a graph illustrating the effect of the Cu content
on the relationship between .DELTA.TS and the heat treatment
temperature after a pre-deformation and a heat treatment of a
hot-dip galvanized steel sheet.
DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0118] A high-ductility steel sheet of the present invention has a
tensile strength TS of not less than 440 MPa, a composite structure
comprising a primary phase containing a ferrite phase and a
secondary phase containing a retained austenite phase with a volume
ratio of not less than 1%, excellent press formability, and
excellent strain age hardenability, which is indicated by a
remarkably increased tensile strength .DELTA.TS of not less than 80
MPa during a heat treatment at a relatively low temperature after
press forming. The term "primary phase" used in the present
invention shall be a structure occupying not less than 50% by a
volume ratio.
[0119] The term "high-ductility steel sheet" used in the present
invention shall mean that a steel sheet has a balance (TS.times.El)
of a tensile strength (TS) and an elongation (El) of not less than
19,000 MPa%.
[0120] In addition, the term ".DELTA.TS" used in the present
invention means an increment in tensile strength between before and
after the heat treatment at a temperature in the range of 150 to
350.degree. C. for a holding time of not less than 30 seconds of a
steel sheet which was subjected to a pre-deformation treatment of a
tensile plastic strain of not less than 5%. That is,
.DELTA.TS=(tensile strength after heat treatment)-(tensile strength
before pre-deformation treatment). The steel sheets of the present
invention shall include hot-rolled steel sheets, cold-rolled steel
sheets and hot-dip galvanized steel sheets.
[0121] All the steel sheets (hot-rolled steel sheets, cold-rolled
steel sheets and hot-dip galvanized steel sheets) having the
above-mentioned structure have high-ductility, excellent press
formability, and excellent strain age hardenability.
[0122] The term "superior strain age hardenability" or the term
"excellent strain age hardenability" used in the present invention
shall mean that, when a steel sheet is subjected to a
pre-deformation treatment of a tensile plastic strain of not less
than 5%, and then, to a heat treatment at a temperature in the
range of 150 to 350.degree. C. for a holding time of not less than
30 seconds, the increment .DELTA.TS in tensile strength between
before and after the heat treatment is not less than 80 MPa,
wherein .DELTA.TS=(tensile strength TS.sub.HT after heat
treatment)-(tensile strength TS before pre-deformation treatment) .
Preferably, the increment .DELTA.TS is not less than 100 MPa. The
heat treatment causes an increase .DELTA.YS in yield stress of not
less than 80 MPa, wherein .DELTA.YS=(yield stress YS.sub.HT after
heat treatment)-(yield stress YS before pre-deformation
treatment).
[0123] In the control of the strain age hardenability, the amount
of prestrain (pre-deformation) plays an important role. The present
inventors investigated the effect of the amount of prestrain on the
subsequent strain age hardenability by assuming possible
deformation types applied to automotive steel sheets. The results
show that the uniaxial equivalent strain (tensile strain) is
generally useful for elucidating the deformation of the steel
sheets except for very deep drawing, that the uniaxial equivalent
strain is mostly more than 5% for actual parts, and that the
strength of the parts exhibit good correspondence to the strength
obtained after a strain aging treatment of a prestrain of 5%. Based
on these findings, a tensile plastic strain of not less than 5% is
employed in the present invention.
[0124] Conventional bake treatment conditions include 170.degree.
C..times.20 minutes as a standard. If precipitation strengthening
by very fine Cu or fine carbide is performed as in the present
invention, the heat treatment temperature must be 150.degree. C. or
more. Under conditions including a temperature exceeding
350.degree. C., on the other hand, the strengthening effect is
saturated, and the steel sheet tends to soften. Heating to a
temperature exceeding 350.degree. C. causes marked occurrence of
thermal strain or temper color. For these reasons, a heat treatment
temperature in the range of 150 to 350.degree. C. is adopted for
strain age hardening in the present invention. The holding time of
the heat treatment temperature should be at least 30 seconds.
Holding a heat treatment temperature in the range of 150 to
350.degree. C. for about 30 seconds permits achievement of
substantially satisfactory strain age hardening. For further
enhanced strain age hardening, the holding time is preferably at
least 60 seconds, and more preferably at least 300 seconds.
[0125] The heat treatment method after the pre-deformation is not
limited in the present invention, and atmospheric heating in a
furnace in general bake treatment, induction heating, non-oxidizing
flame heating, laser heating, and plasma heating are suitably
applicable. So-called hot pressing for pressing a heated steel
sheet is also very effective means in the present invention.
[0126] Next, the hot-rolled steel sheet, the cold-rolled steel
sheet, and the hot-dip galvanized steel sheet in the present
invention will be described individually.
[0127] (1) Hot-rolled Steel Sheet
[0128] The hot-rolled steel sheet of the present invention will now
be described.
[0129] The hot-rolled steel sheet of the present invention has a
composite structure comprising a ferrite primary phase and a
secondary phase containing a retained austenite phase having a
volume ratio of not less than 1% of the entire structure. As
described above, a hot-rolled steel sheet having such a composite
structure exhibits high ductility, high strength-ductility balance
(TS.times.El), and excellent press formability.
[0130] Ferrite primary phase is preferably present in a volume
ratio of not less than 50%. With a ferrite phase of less than 50%,
it is difficult to keep high ductility, resulting in lower press
formability. When further enhanced ductility is required, the
volume ratio of the ferrite phase is preferably not less than 80%.
For the purpose of making full use of advantages of the composite
structure, the ferrite phase is preferably not more than 98%.
[0131] In the present invention, steel must contain retained
austenite phase as the secondary phase in a volume ratio of not
less than 1% of the entire structure. With a retained austenite
phase of less than 1%, high elongation (El) cannot be obtained. To
obtain higher elongation (El), the retained austenite phase content
is preferably not less than 2% and more preferably not less than
3%.
[0132] The secondary phase may be a single retained austenite phase
having a volume ratio of not less than 1%, or may be a mixture of a
retained austenite phase of a volume ratio of not less than 1% and
another phase, i.e., a pearlite phase, a bainite phase, and/or a
martensite phase.
[0133] The reasons for limiting the composition of the hot-rolled
steel sheet of the present invention will now, be described. The
weight percent in the composition will hereafter be denoted simply
as %.
[0134] C: 0.05 to 0.20%
[0135] C is an element, which improves strength of a steel sheet
and promotes the formation of a composite structure of ferrite and
retained austenite, and is preferably contained in an amount of not
less than 0.05% for forming the composite structure according to
the present invention. A C content exceeding 0.20% causes an
increase in proportions of carbides in steel, resulting in a
decrease in ductility, and hence a decrease in press formability. A
more serious problem is that a C content exceeding 0.20% leads to
significant deterioration of spot weldability and arc weldability.
For these reasons, the C content is limited within the range of
0.05 to 0.20% in the present invention. From the viewpoint of
formability, the C content is preferably not more than 0.18%.
[0136] Si: 1.0 to 3.0%
[0137] Si is a useful strengthening element, which improves the
strength of a steel sheet without a marked decrease in ductility of
the steel sheet. In addition, Si is necessary for forming a
retained austenite phase. To obtain these effects, Si is preferably
contained in an amount of not less than 1.0% and more preferably
not less than 1.2%. An Si content exceeding 3.0% leads to
deterioration of press formability and degrades the surface
quality. The Si content is therefore limited within the range of
1.0 to 3.0%.
[0138] Mn: not more than 3.0%
[0139] Mn is a useful element, which strengthens steel and prevents
hot cracking caused by S, and is therefore contained in an amount
according to the S content. These effects are particularly
remarkable at an Mn content of not less than 0.5%. On the other
hand, an Mn content exceeding 3.0% results in deterioration of
press formability and weldability. The Mn content is therefore
limited to not more than 3.0% in the present invention. More
preferably, the Mn content is not less than 1.0%.
[0140] P: not more than 0.10%
[0141] P strengthens steel, and may be contained in an amount
necessary for a desired strength. From the viewpoint of increasing
the strength, P is preferably contained in an amount of not less
than 0.005%. On the other hand, a P content exceeding 0.10% results
in deterioration of press formability. The P content is therefore
limited to not more than 0.10%. When superior press formability is
required, the P content is preferably not more than 0.08%.
[0142] S: not more than 0.02%
[0143] S is an element, which is present as inclusions in a steel
sheet and causes deterioration of ductility, formability, and
particularly stretch flanging formability of the steel sheet, and
it should be the lowest possible. A reduced S content of not more
than 0.02% does not exert much adverse effect and therefore, the S
content is limited to up to 0.02% in the present invention. When
more excellent stretch flanging formability is required, the S
content is preferably not more than 0.010%.
[0144] Al: not more than 0.30%
[0145] Al is a useful element, which is added as a deoxidizing
element to steel, and improves cleanliness of steel. In addition,
Al facilitates the formation of the retained austenite. These
effects are particularly remarkable at an Al content of not less
than 0.01%. The Al content exceeding 0.30% cannot give further
effects, but causes deterioration of press formability. The Al
content is therefore limited to not more than 0.30%. Preferably,
the Al content is not more than 0.10%. The present invention does
not exclude a steelmaking process based on deoxidation using a
deoxidizer other than Al. For example, Ti deoxidation or Si
deoxidation may be employed, and steel sheets produced by such
deoxidation methods are also included in the scope of the present
invention. In this case, addition of Ca or REM to molten steel does
not impair the features of the steel sheet of the present invention
at all.
[0146] N: not less than 0.02%
[0147] N is an element, which increases the strength of a steel
sheet through solid solution strengthening or strain age hardening,
and is preferably contained in an amount of not less than 0.0010%
to obtain these effects. However, an N content exceeding 0.02%
causes an increase in the content of nitrides in the steel sheet,
which causes serious deterioration of ductility, and thus, of press
formability of the steel sheet. The N content is therefore limited
to not more than 0.02%. When further improvement in press
formability is required, the N content is preferably not more than
0.01%, and more preferably less than 0.0050%.
[0148] Cu: 0.5 to 3.0%
[0149] Cu is an element, which remarkably increases strain age
hardening of a steel sheet (increase in strength after
pre-deformation/heat treatment), and thus is most important in the
present invention. With a Cu content of less than 0.5%, an
increment .DELTA.TS in tensile strength exceeding 80 MPa cannot be
obtained by changing the pre-determination/heat treatment
conditions. With a Cu content exceeding 3.0%, the effect is
saturated so that an effect corresponding to the content cannot be
expected, leading to unfavorable economic effects. Furthermore,
deterioration of press formability occurs, and the surface quality
of the steel sheet is degraded. The Cu content is therefore limited
within a range of 0.5 to 3.0%. In order to simultaneously achieve a
higher .DELTA.TS and excellent press formability, the Cu content is
preferably within a range of 1.0 to 2.5%.
[0150] The hot-rolled steel sheet of the present invention
containing Cu preferably further contains, in weight percent, at
least one of the following Groups A to C:
[0151] Group A: Ni: not more than 2.0%;
[0152] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0153] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0154] Group A: Ni: not more than 2.0%
[0155] Group A: Ni is effective for preventing the formation of
surface defects on the steel sheet surface containing Cu, and may
be added as required. The Ni content is preferably about a half the
Cu content, i.e., in the range of about 30 to about 80% of the Cu
content. An Ni content exceeding 2.0% cannot give further
enhancement in the effect because saturation of the effect, leading
to economic disadvantages, and causes deterioration of press
formability. For these reasons, the Ni content is preferably
limited to not more than 2.0%.
[0156] Group B: at least one of Cr and Mo: not more than 2.0% in
total
[0157] Group B: Both Cr and Mo, as well as Mn, strengthen the steel
sheet and at least one thereof can be contained as required. This
effect is particularly remarkable at a Cr content of not less than
0.1% and at an Mo content of not less than 0.1%. It is therefore
preferable to contain at least one of Cr: not less than 0.1% and
Mo: not less than 0.1%. If at least one of Cr and Mo are contained
in a total amount exceeding 2.0%, press formability is impaired. It
is therefore preferable to limit the total content of Cr and Mo to
not more than 2.0%.
[0158] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total
[0159] Group C: Nb, Ti, and V are carbide-forming elements and
effectively increase the strength by fine dispersion of carbides,
and can be selected and contained as required. This effect can be
achieved at an Nb content of not less than 0.01%, a Ti content of
not less than 0.01%, and a V content of not less than 0.01%.
However, a total content of Nb, Ti, and V exceeding 0.2% causes
deterioration of press formability. Thus, the total content of Nb,
Ti, and V is preferably limited to not more than 0.2%.
[0160] In the present invention, in place of the aforementioned Cu
or at least one of the above-mentioned Groups A to C, at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0%, and W: 0.05 to 2.0% may be contained in an amount of not more
than 2.0% in total, and at least one selected from the group
consisting of Nb, Ti, and V may be further contained in an amount
of not more than 2.0% in total.
[0161] At least one selected from the group consisting of Mo: 0.05
to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not
more than 2.0% in total
[0162] Mo, Cr, and W are elements, which remarkably increase strain
age hardening (increase in strength after pre-deformation and heat
treatment) of a steel sheet, and are one of the most important
elements in the present invention. That is, in the present
invention, a hot-rolled steel sheet having a composite structure
containing ferrite as a primary phase and a secondary phase of
retained austenite and containing at least one of Mo, Cr, and W,
causes strain-induced transformation of the retained austenite into
martensite when a prestrain of not less than 5% and a
low-temperature heat treatment are applied to the hot-rolled steel
sheet, and strain-induced fine precipitation of fine carbides at a
low temperature occurs in the strain-induced transformed
martensite, resulting in an increase in tensile strength .DELTA.TS
of not less than 80 MPa. With a content of at least one of Mo, Cr,
and W of less than 0.05%, changing the steel sheet structure and
pre-deformation and heat treatment conditions does not cause an
increase in tensile strength .DELTA.TS of not less than 80 MPa. On
the other hand, a content of at least one of Mo, Cr, and W
exceeding 2.0% does not give a corresponding effect because of
saturation of the effect, leading to economic disadvantages, and
causes deterioration of press formability. The contents of Mo, Cr,
and W are each preferably limited within the range of 0.05 to 2.0%.
From the viewpoint of press formability, the total content of Mo,
Cr and/or W is more preferably limited to not more than 2.0%.
[0163] At least one of Nb, Ti, and V, in a total amount of not more
than 2.0%
[0164] Nb, Ti, and V are carbide-forming elements, and can be added
as required. Containing at least one of Nb, Ti, and V, in addition
to at least one of Mo, Cr, and W, and forming a composite structure
containing a ferrite primary phase and a secondary phase of
retained austenite form fine carbides in the strain-induced
transformed martensite and cause strain-induced precipitation at
low temperature, resulting in an increase in tensile strength
.DELTA.TS of not less than 80 MPa. In order to obtain these
effects, an Nb content is preferably not less than 0.01%, a Ti
content is preferably not less than 0.01%, and a V content is
preferably not less than 0.01%, and at least one of Nb, Ti, and V
can be added as required. However, a total content exceeding 2.0%
causes deterioration of press formability. Thus, the total content
of Nb, Ti, and V is preferably limited to not more than 2.0%.
[0165] Apart from the above-mentioned elements, at least one of Ca:
not less than 0.1% and REM: not less than 0.1% may be contained. Ca
and REM are elements contributing to improvement in stretch
flanging property through conformational control of inclusions. If
the Ca content exceeds 0.1% or the REM content exceeds 0.1%,
however, there would be a decrease in cleanliness, and a decrease
in ductility.
[0166] The balance of the composition of the steel sheet is Fe and
incidental impurities. Allowable incidental impurities are Sb: not
more than 0.01%, Sn: not more than 0.1%, Zn: not more than 0.01%,
Co: not more than 0.1%, Zr: not more than 0.1%, and B: not more
than 0.1%.
[0167] A method for manufacturing the hot-rolled steel sheet of the
present invention will now be described.
[0168] The hot-rolled steel sheet of the present invention is made
by hot-rolling a steel slab having a composition within the ranges
described above into a prescribed thickness.
[0169] While the steel slab used is preferably manufactured by a
continuous casting process to prevent macro-segregation of the
constituents, it may be manufactured by an ingot casting process or
a thin-slab casting process. A conventional process employed in
this embodiment includes the steps of manufacturing a steel slab,
cooling the steel slab to room temperature, and reheating the slab.
Alternatively, an energy-saving process also is applicable without
problem in the present invention. For example, a hot steel slab is
charged into a heating furnace without cooling to room temperature,
or directly rolled immediately after short temperature
holding(direct-hot-charge rolling or direct rolling).
[0170] The reheating temperature SRT of the material (steel slab)
is not limited and is preferably not less than 900.degree. C.
[0171] Slab reheating temperature: not less than 900.degree. C.
[0172] The slab reheating temperature (SRT) is preferably the
lowest possible with a view to prevent surface defects caused by Cu
when the material contains Cu. However, with a reheating
temperature of less than 900.degree. C., there is an increase in
the rolling load, thus increasing the risk of occurrence of a
trouble during hot rolling. Considering the increase in scale loss
caused along with accelerated oxidation, the slab reheating
temperature is preferably not more than 1,300.degree. C.
[0173] From the viewpoint of reducing the slab reheating
temperature and preventing occurrence of troubles during hot
rolling, use of a so-called sheet bar heater heating a sheet bar is
of course an effective method.
[0174] The reheated steel slab is then hot-rolled into a hot-rolled
sheet. In the present invention, a finish rolling condition is
particularly important, and the hot rolling is preferably performed
at a finish rolling end temperature (FDT) in the range of 780 to
980.degree. C.
[0175] At the FDT of less than 780.degree. C., a deformed structure
remains in the steel sheet to cause deterioration of ductility. On
the other hand, an FDT exceeding 980.degree. C. coarsens the
structure, leading to a decrease in formability due to delay of
ferrite transformation. Thus, the FDT is preferably in the range of
780 to 980.degree. C.
[0176] After the completion of finish rolling, a forced cooling
treatment is applied. In the present invention, a forced cooling
condition is particularly important. In the present invention,
within 2 seconds after the completion of finish rolling, a forced
cooling is preferably carried out at a cooling rate of not less
than 50.degree. C./second to a temperature in the range of 620 to
780.degree. C. With a cooling start time exceeding 2 seconds, the
structure coarsens and ferrite transformation is delayed, resulting
in poor press formability. The cooling start time after the
completion of finish rolling is preferably limited to within 2
seconds.
[0177] With a cooling rate of less than 50.degree. C./second after
the completion of finish rolling, and ferrite transformation
undesirably starts during the forced cooling, ferrite
transformation does not appropriately occur in a subsequent
isothermal holding treatment or slow cooling treatment, thus
resulting in a decreased press formability. Accordingly, the
cooling rate is preferably limited to not less than 50.degree.
C./second. However, with a cooling rate exceeding 300.degree.
C./second, degradation of the steel sheet shape is concerned. Thus,
the upper limit of the cooling rate is preferably 300.degree.
C./second.
[0178] In addition, in the present invention, the steel sheet is
preferably cooled to the vicinity of a nose of a free or
pro-eutectoid ferrite temperature region of 620 to 780.degree. C.
by the above-mentioned forced cooling. At a cooling stop
temperature of less than 620.degree. C. of the forced cooling, free
ferrite is not generated, but pearlite is generated. At a cooling
stop temperature exceeding 780.degree. C., a decrease in
concentration of carbon into austenite decreases with a decrease in
the generation of free ferrite. The cooling stop temperature of
forced cooling is more preferably in the range of 650 to
750.degree. C.
[0179] After the forced cooling to the vicinity of a nose of free
ferrite temperature region of 620 to 780.degree. C., an isothermal
holding treatment for 1 to 10 seconds within the above-mentioned
temperature region or a slow cooling treatment at a cooling rate of
not more than 20.degree. C./second is preferably performed.
[0180] By the isothermal holding treatment for a short period of
time within this temperature region (620 to 780.degree. C.) or the
slow cooling treatment for a short period of time within the
above-mentioned temperature region, a desired amount of free
ferrite can be formed.
[0181] For achieving the concentration of carbon into the austenite
along with ferrite transformation, the isothermal holding treatment
or slow cooling treatment is more preferably performed within a
temperature region of 620.degree. C. to 750.degree. C.
[0182] A holding time of the isothermal treatment or a time
required for the slow cooling treatment of less than 1 second
causes insufficient concentration of carbon into the austenite. On
the other hand, a time exceeding 10 seconds causes pearlite
transformation.
[0183] A cooling rate of the slow cooling treatment exceeding
20.degree. C./second causes insufficient concentration of carbon
into the austenite.
[0184] After the isothermal holding treatment or slow cooling
treatment, the rolled sheet is preferably cooled again to a
temperature of 300 to 500.degree. C. at a cooling rate of not less
than 50.degree. C./second, and then coiled. That is, the rolled
sheet is preferably coiled at a coiling temperature (CT) of 300 to
500.degree. C.
[0185] After the isothermal holding treatment or slow cooling
treatment, the rolled sheet is cooled to a temperature of 300 to
500.degree. C. Also, the cooling rate of this treatment is
preferably not less than 50.degree. C./second. With the cooling
rate of less than 50.degree. C./second, pearlite transformation
occurs and ductility is decreased. The cooling rate is more
preferably within the range of 50 to 200.degree. C./second.
[0186] With a coiling temperature CT of less than 300.degree. C.,
the secondary phase contains martensite. On the other hand, with
the coiling temperature exceeding 500.degree. C., the secondary
phase contains pearlite. Thus, the coiling temperature CT is
preferably within a range of 300 to 500.degree. C.
[0187] In the present invention, all or part of finish rolling may
be lubrication rolling to reduce the rolling load during hot
rolling. Application of lubrication rolling is effective also from
the viewpoint of achieving a uniform steel sheet shape and uniform
material quality. The frictional coefficient on the lubrication
rolling is preferably in the range of 0.25 to 0.10. A continuous
rolling process is preferable one,in which neighboring sheet bars
can be connected to each other to perform finish rolling
continuously. Application of the continuous rolling process is
desirable also from the viewpoint of operational stability of hot
rolling.
[0188] After the completion of hot rolling, temper rolling of not
more than 10% may be applied for adjustment such as shape
correction or surface roughness control.
[0189] The hot-rolled steel sheet of the invention may be used as a
steel sheet for processing and as a steel sheet for surface
treatments. Surface treatments include galvanizing (including
alloying), tin-plating and enameling. After annealing or
galvanizing, the hot-rolled steel sheet of the present invention
may be subjected to a special treatment to improve activity to
chemical treatment, weldability, press formability, and corrosion
resistance.
[0190] (2) Cold-rolled Steel Sheet
[0191] A cold-rolled steel sheet of the present invention will now
be described.
[0192] The cold-rolled steel sheet of the present invention has a
composite structure comprising a ferrite primary phase and a
secondary phase containing retained austenite having a volume ratio
of not less than 1% of the entire structure. As described above, a
cold-rolled steel sheet having such a composite structure exhibits
high elongation (El), high strength/elongation balance
(TS.times.El), and excellent press formability.
[0193] The volume ratio of the ferrite primary phase contained in
the composite structure is preferably not less than 50%. With a
ferrite phase content of less than 50%, it is difficult to keep
high ductility, resulting in poor press formability. When further
enhanced ductility is required, the volume ratio of the ferrite
phase is preferably not less than 80%. For the purpose of making
full use of advantages of the composite structure, the ferrite
phase is preferably not more than 98%.
[0194] In the present invention, the steel sheet must contain a
retained austenite phase as the secondary phase in a volume ratio
of not less than 1% of the entire structure. With a retained
austenite phase content of less than 1%, it is impossible to obtain
high elongation (El). To obtain higher elongation (El), the
retained austenite phase is preferably contained in a volume ratio
of not less than 2%, more preferably, not less than 3%.
[0195] The secondary phase may be a single retained austenite phase
having a volume ratio of not less than 1%, or may be a mixture of a
retained austenite phase of a volume ratio of not less than 1% and
an auxiliary (another) phase comprising a pearlite phase, a bainite
phase, and/or a martensite phase.
[0196] The reasons for limiting the composition of the cold-rolled
steel sheet of the present invention will now be described. The
weight percent in the composition will simply be denoted
hereinafter as %.
[0197] C: not more than 0.20%
[0198] C is an element, which improves strength of a steel sheet
and promotes the formation of a composite structure of a ferrite
phase and a retained austenite phase, and is preferably contained
in an amount of not less than 0.01% from the viewpoint of forming
the retained austenite phase in the present invention. A C content
is more preferably not less than 0.05%. A C content exceeding
0.20%, however, causes an increase in amount of carbides in the
steel, resulting in a decrease in ductility, and hence a decrease
in press formability. A more serious problem is that a C content
exceeding 0.20% leads to remarkable deterioration of spot
weldability and arc weldability. For these reasons, in the present
invention, the C content is limited to not more than 0.20%. From
the viewpoint of formability, the C content is preferably not more
than 0.18%.
[0199] Si: not more than 2.0%
[0200] Si is a useful strengthening element, which improves
strength of a steel sheet without a marked decrease in ductility of
the steel sheet and facilitates the formation of a residual
austenite phase. The Si content is preferably not less than 0.1%.
An Si content exceeding 2.0%, however, leads to deterioration of
press formability and degrades the surface quality. The Si content
is, therefore, limited to not more than 2.0%.
[0201] Mn: not more than 3.0%
[0202] Mn is a useful element, which strengthens the steel and
prevents hot cracking caused by S, and is therefore contained in an
amount according to the S content. These effects are particularly
remarkable at an Mn content of not less than 0.5%. However, an Mn
content exceeding 3.0% results in deterioration of press
formability and weldability. The Mn content is, therefore, limited
to not more than 3.0% in the present invention. More preferably,
the Mn content is not less than 1.0%.
[0203] P: not more than 0.10%
[0204] P strengthens the steel, and may be contained in an amount
of preferably not less than 0.005%., according to a desired
strength. However, an excess P content causes deterioration of
press formability. The P content is, therefore, limited to not more
than 0.10%. When more excellent press formability is required, the
P content is preferably not more than 0.08%.
[0205] S: not more than 0.02%
[0206] S is an element, which is present as inclusions in steel and
causes deterioration of ductility, formability, and particularly
stretch flanging formability of a steel sheet, and it should be the
lowest possible. However, an S content reduced to not more than
0.02% does not exert much adverse effect. Thus, the S content is
limited to not more than 0.02% in the present invention. When
superior stretch flanging formability is required, the S content is
preferably not more than 0.010%.
[0207] Al: not more than 0.30%
[0208] Al is a deoxidizing element of steel, and is useful for
improving cleanliness of the steel. In addition, Al is effective
for the formation of the retained austenite. In order to obtain
these effects, the Al content is preferably not less than 0.01%.
However, an Al content exceeding 0.30% cannot give further enhanced
deoxidizing effects, and causes deterioration of press formability.
The Al content is, therefore, limited to not more than 0.30%. The
invention also includes a steel making process using other
deoxidizers, for example, Ti or Si, and steel sheets produced by
such deoxidation methods are also included in the scope of the
invention. In this case, addition of Ca or REM to molten steel does
not impair the features of the steel sheet of the invention at all.
Of course, steel sheets containing Ca or REM are included within
the scope of the invention.
[0209] N: not more than 0.02%
[0210] N is an element, which increases strength of a steel sheet
through solid solution strengthening or strain age hardening, and
is preferably contained in an amount of not more than 0.001%.
However, an N content exceeding 0.02% causes an increase in nitride
content in the steel sheet, whereby ductility and press formability
of the steel sheet are seriously deteriorated. The N content is
therefore limited to not more than 0.02%. When further improvement
of press formability is required, the N content is preferably not
more than 0.01%.
[0211] Cu: 0.5 to 3.0%
[0212] Cu is an element, which remarkably increases strain age
hardening of a steel sheet (increase in strength after
pre-deformation/heat treatment), and is one of the most important
elements in the present invention. With a Cu content of less than
0.5%, an increase in tensile strength .DELTA.TS exceeding 80 MPa
cannot be obtained by changing the pre-deformation/heat treatment
conditions. In the present invention, therefore, Cu should be
contained in an amount of not less than 0.5%. With a Cu content
exceeding 3.0%, however, the effect is saturated, leading to
unfavorable economic effects. Furthermore, deterioration of press
formability occurs, and the surface quality of the steel sheet is
degraded. The Cu content is, therefore, limited within the range of
0.5 to 3.0%. In order to simultaneously achieve a higher .DELTA.TS
and excellent press formability, the Cu content is preferably
within the range of 1.0 to 2.5%.
[0213] In the present invention, the above-mentioned composition
containing Cu preferably further contains, in weight percent, at
least one of the following Groups A to C:
[0214] Group A: Ni: not more than 2.0%;
[0215] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0216] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0217] Group A: Ni: not more than 2.0%
[0218] Group A: Ni is an element effective for preventing surface
defects produced by Cu contained in the steel sheet, and may be
contained as required. The Ni content depends on the Cu content,
and is preferably about a half the Cu content, more specifically,
within the range of about 30 to about 80% of the Cu content. An Ni
content exceeding 2.0% cannot give further enhancement in the
effect because of saturation of the effect, leading to economic
disadvantages, and causes deterioration of press formability. For
these reasons, the Ni content is preferably limited to not more
than 2.0%.
[0219] Group B: at least one of Cr and Mo: not more than 2.0% in
total
[0220] Group B: Both Cr and Mo, as well as Mn, strengthen the steel
sheet and may be contained as required preferably in an amount of
not less than 0.1% for Cr and not less than 0.1% for Mo. If at
least one of Cr and Mo are contained in an amount exceeding 2.0% in
total, press formability is impaired. It is therefore preferable to
limit the total content of Cr and Mo forming Group B to not more
than 2.0%.
[0221] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total
[0222] Group C: Nb, Ti, and V are elements, which effectively form
fine dispersion of carbides contributing to an increase in
strength. Therefore, Nb, Ti, and V can be selected and contained as
required preferably in an amount of not less than 0.01% for Nb, in
an amount of not less than 0.01% for Ti and in an amount of not
less than 0.01% for V. If the total content of at least one of Nb,
Ti, and V exceeds 0.2%, the press formability is impaired. Thus,
the total content of Nb, Ti and/or V is preferably limited to not
more than 0.2%.
[0223] In the present invention, in place of the aforementioned Cu,
at least one selected from the group consisting of Mo: 0.05 to
2.0%, Cr: 0.05 to 2.0%, and W: 0.05 to 2.0% may be contained in an
amount of not more than 2.0% in total.
[0224] At least one selected from the group consisting of Mo: 0.05
to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not
more than 2.0% in total
[0225] In the present invention, all of Mo, Cr, and W, as well as
Cu, are the most important elements, which remarkably increase
strain age hardening of the steel sheet, and can be selected and
contained. When a steel sheet containing at least one of Mo, Cr,
and W and having a composite structure of a ferrite phase and a
phase containing retained austenite is subjected to a prestrain
(pre-deformation) of not less than 5% and a low-temperature heat
treatment (heat treatment), the retained austenite is changed into
martensite by strain-induced transformation. Then, the formation of
fine carbide precipitation in the martensite is induced by the
strain, resulting in an increase in tensile strength .DELTA.TS of
not less than 80 MPa. With a content of each of these elements of
less than 0.05%, changing pre-deformation/heat treatment conditions
does not give an increase in tensile strength .DELTA.TS of at least
80 MPa. If the content of each of these elements exceeds 2.0%, a
further enhanced effect corresponding to the content cannot be
expected as a result of saturation of the effect, leading to
economic disadvantages, and this results in deterioration of press
formability. The contents of Mo, Cr, and W are therefore limited
within the range of 0.05 to 2.0% for Mo, 0.05 to 2.0% for Cr, and
0.05 to 2.0% for W. From the viewpoint of press formability, the
total content of Mo, Cr, and W is limited to not more than
2.0%.
[0226] In the present invention, at least one selected from the
group consisting of Mo, Cr, and W is preferably contained and
further, at least one of Nb, Ti, and V are preferably contained not
more than 2.0% in total.
[0227] At least one of Nb, Ti, and V, in a total amount of not more
than 2.0%:
[0228] Nb, Ti, and V are elements forming carbides, and can be
selected and contained as required, when at least one of Mo, Cr,
and W is added. When the steel composition contains at least one of
Mo, Cr, and W and has a composite structure containing a ferrite
phase and a retained austenite phase, and contains at least one of
Nb, Ti, and V, the retained austenite is transformed into
martensite by strain-induced transformation during the
pre-deformation/heat treatment. Then, fine carbide precipitation is
induced by the strain in the martensite, thus resulting in an
increase in tensile strength .DELTA.TS of not less than 80 MPa.
This effect is particularly remarkable preferably at a Nb content
of not less than 0.01%, at a Ti content of not less than 0.01%, and
at a V content of not less than 0.01%. However, a total content of
Nb, Ti, and V exceeding 2.0% causes deterioration of press
formability. Thus, the total content of Nb, Ti and/or V is
preferably limited to not more than 2.0%.
[0229] Although no particular restriction is imposed, apart from
the above-mentioned constituents, the composition may contain B:
not more than 0.1%, Zr: not more than 0.1%, Ca: not more than 0.1%,
and REM: not more than 0.1% without any problem.
[0230] The balance of the composition of the steel is Fe and
incidental impurities. Allowable incidental impurities include Sb:
not more than 0.01%, Sn: not more than 0.1%, Zn: not more than
0.01%, and Co: not more than 0.1%.
[0231] The method for manufacturing the cold-rolled steel sheet of
the present invention will now be described.
[0232] The cold-rolled steel sheet of the present invention is
manufactured through a hot rolling step of hot-rolling a steel slab
having the composition within the aforementioned ranges into a
hot-rolled steel sheet, a cold rolling step, of cold-rolling the
hot-rolled steel sheet into a cold-rolled steel sheet, and a
recrystallization annealing step of recrystallization-annealing the
cold-rolled steel sheet to form a cold-rolled annealed steel
sheet.
[0233] Although the steel slab used is preferably manufactured by a
continuous casting process to prevent macrosegregation of the
constituents, it may be manufactured by an ingot casting process or
a thin-slab continuous casting process. A conventional process
employed in this embodiment includes the steps of manufacturing a
steel slab, cooling the steel slab to room temperature, and
reheating the slab. Alternatively, an energy-saving process is
applicable without problem in the present invention. For example, a
hot steel slab is charged into a reheating furnace without cooling
to room temperature, or directly rolled immediately after short
temperature holding (direct-feed rolling or direct rolling).
[0234] The steel slab having the above-mentioned composition is
reheated and hot-rolled to make a hot-rolled steel sheet. No
particular problem is encountered as to conventionally known
conditions so far as such conditions permit manufacture of a
hot-rolled steel sheet having a desired thickness in the hot
rolling step. Preferable conditions for hot rolling are as
follows:
[0235] Slab reheating temperature: not less than 900.degree. C.
[0236] The slab reheating temperature is preferably the lowest
possible with a view to prevent surface defects caused by Cu when
the composition contains Cu. However, with a reheating temperature
of less than 900.degree. C., the rolling load increases, thus
increasing the risk of occurrence of a trouble during hot rolling.
In view of an increase in scale loss caused by facilitated
oxidation, the slab reheating temperature is preferably not more
than 1,300.degree. C.
[0237] From the viewpoint of reducing the slab reheating
temperature and preventing occurrence of troubles during hot
rolling, use of a so-called sheet bar heater, which heats a sheet
bar, is effective.
[0238] Finish rolling end temperature: not less than 700.degree.
C.
[0239] At a finish rolling end temperature (FDT) of not less than
700.degree. C., it is possible to obtain a uniform hot-rolled
mother sheet structure which can give an excellent formability
after cold rolling and recrystallization annealing. A finish
rolling end temperature of less than 700.degree. C. leads to a
non-uniform structure of the hot-rolled mother sheet and a higher
rolling load during hot rolling, thus increasing the risk of
occurrence of troubles during hot rolling. Thus, the FDT for the
hot rolling step is preferably not less than 700.degree. C.
[0240] Coiling temperature: not more than 800.degree. C.
[0241] The coiling temperature is preferably not more than
800.degree. C., and more preferably not less than 200.degree. C. A
coiling temperature exceeding 800.degree. C. tends to cause a
decrease in yield as a result of an increased scale loss. With a
coiling temperature of less than 200.degree. C., the steel sheet
shape is seriously impaired, and there is an increasing risk of
occurrence of inconveniences in practical use.
[0242] In the hot rolling step in the present invention, as
described above, it is desirable to reheat the slab to a
temperature of not less than 900.degree. C., hot-roll the reheated
slab at a finish rolling end temperature of not less than
700.degree. C., and coil the hot-rolled steel sheet at a coiling
temperature of not more than 8000C and preferably not less than
200.degree. C.
[0243] In the hot rolling step in the present invention, all or
part of finish rolling may be lubrication rolling, which reduces
the rolling load during the hot rolling. The lubrication rolling is
effective also from the viewpoint of achieving a uniform steel
sheet shape and a uniform material quality. The frictional
coefficient on the lubrication rolling is preferably within a range
of 0.25 to 0.10. It is desirable to connect neighboring sheet bars
to each other to perform a continuous finish rolling process.
Application of the continuous rolling process is desirable also
from the viewpoint of operational stability of hot rolling.
[0244] Then, a cold rolling step is conducted for the hot-rolled
steel sheet. In the cold rolling step, the hot-rolled steel sheet
is cold-rolled into a cold-rolled steel sheet. Any cold rolling
conditions may be used so far as such conditions permit production
of cold-rolled steel sheets with desired dimensions and shape, and
no particular restriction is imposed. The reduction in cold rolling
is preferably not less than 40%. With a reduction of less than 40%,
uniform recrystallization barely occurs during the subsequent
recrystallization-annealing step.
[0245] Then, the cold-rolled steel sheet is subjected to the
recrystallization annealing step to convert the sheet into a
cold-rolled annealed steel sheet. The recrystallization annealing
is preferably carried out on a continuous annealing line. In the
present invention, the recrystallization annealing is a heat
treatment which includes heating and soaking the cold-rolled sheet
in the dual phase region of ferrite and austenite in the
temperature range between the A.sub.C1 transformation point and the
A.sub.C3 transformation point, cooling the sheet, and retaining the
sheet at a temperature in the range of 300 to 500.degree. C. for 30
to 1,200 seconds.
[0246] The heating and soaking temperature for recrystallization
annealing is preferably within the dual phase region in the
temperature range between the A.sub.C1 transformation point and the
AC.sub.3 transformation point. The heating and soaking temperature
of less than the A.sub.C1 transformation point leads to the
formation a single ferrite phase. On the other hand, a high
temperature exceeding AC.sub.3 transformation point results in
coarsening of crystal grains, the formation of a single austenite
phase, and a serious deterioration of press formability.
[0247] After the heating and soaking treatment, the sheet was
cooled from the heating and soaking temperature and retained at a
temperature in the range of 300 to 500.degree. C. for 30 to 1,200
seconds. The heating and soaking treatment and the subsequent
retaining treatment facilitates the formation of a retained
austenite phase of not less than 1%. When the temperature for the
retaining treatment is less than 300.degree. C., the composite
structure of ferrite and martensite is formed. On the other hand, a
temperature range exceeding 500.degree. C. leads to a
ferrite/bainite composite structure or a ferrite/pearlite composite
structure. In these cases, the retained austenite is barely
formed.
[0248] In addition, a retention time of less than 30 seconds in the
temperature range of 300 to 500.degree. C. cannot lead to the
formation of the retained austenite structure. Also, the retention
time exceeding 1,200 seconds cannot lead to the formation of the
retained austenite structure, but leads to a ferrite/bainite
composite structure. Therefore, the retention time in the
temperature region of 300 to 500.degree. C. is preferably in the
range of 30 to 1,200 seconds.
[0249] By the recrystallization annealing, a composite structure of
a ferrite phase and a retained austenite phase is formed, whereby a
high .DELTA.TS can be obtained together with high ductility.
[0250] After the hot rolling, temper rolling with a reduction rate
of not more than 10% may be applied for adjustments and other shape
correction and, surface roughness control.
[0251] The cold-rolled steel sheet of the invention may be used as
a steel sheet for processing and as a steel sheet for
surface-treating. Surface treatments include galvanizing (including
alloying), tin-plating and enameling. After galvanizing, the
cold-rolled steel sheet of the present invention may be subjected
to a special treatment to improve activity to chemical treatment,
weldability, press formability, and corrosion resistance.
[0252] (3) Hot-dip Galvanized Steel Sheet
[0253] The hot-dip galvanized steel sheet of the present invention
will now be described.
[0254] The hot-dip galvanized steel sheet of the present invention
has a composite structure comprising a primary phase consisting of
a ferrite phase and a tempered martensite phase and a secondary
phase containing retained austenite phase in a volume ratio of not
less than 2%.
[0255] Note that the term "tempered martensite phase" in the
present invention means a phase produced by heating a lath
martensite. That is, the tempered martensite phase still maintains
a fine internal structure of the lath martensite, after the heating
(tempering). Furthermore, the tempered martensite phase is softened
by heating (tempering), has enhanced deformability as compared with
martensite, and is effective for improving ductility of the steel
sheet. Note that the term "lath martensite" means martensite
consisting of bundles of thin long platelike martensite crystals,
which can be observed with an electron microscope.
[0256] In the hot-dip galvanized steel sheet of the present
invention, the total volume ratio of the ferrite phase and the
tempered martensite phase functioning as the primary phase is
preferably not less than 50%. With a total volume ratio of the
ferrite phase and the tempered phase of less than 50%, it is
difficult to secure high ductility and press formability is
decreased. When further enhanced ductility is required, the total
volume ratio of the ferrite phase and the martensite phase
functioning as the primary phase is preferably not less than 80%.
For the purpose of making full use of advantages of the composite
structure, the total of the ferrite phase and the tempered
martensite phase is preferably not more than 98%. The ferrite phase
constituting the primary phase preferably occupies not less than
30% by volume of the entire structure, and the tempered martensite
phase preferably occupies not less than 20% by volume of the entire
structure. With a volume ratio of the ferrite phase of less than
30%, or with a volume ratio of the tempered martensite phase of
less than 20%, the ductility will not be remarkably enhanced.
[0257] The hot-dip galvanized steel sheet of the present invention
contains a retained austenite phase as a secondary phase with a
volume ratio of not less than 1% of the entire structure. With a
content of the retained austenite phase of less than 1%, high
elongation (El) cannot be obtained. In order to obtain higher
elongation (El), the retained austenite phase is preferably
contained not less than 2% and more preferably not less than 3%.
The secondary phase may be a single retained austenite phase having
a volume ratio of not less than 1%, or may be a mixture of a
retained austenite phase of a volume ratio of not less than 1% and
an auxiliary (other) phase, for example, a pearlite phase, a
bainite phase, and/or a martensite phase.
[0258] The reasons for limiting the composition of the hot-dip
galvanized steel sheet of the present invention will now be
described.
[0259] C: not more than 0.20%
[0260] C is an element, which improves the strength of a steel
sheet and promotes the formation of a composite structure of a
primary phase comprising ferrite and tempered martensite and a
secondary phase containing retained austenite. In the present
invention, from the viewpoint of formation of the composite
structure, C is preferably contained in an amount of not less than
0.01%. A C content exceeding 0.20% causes an increase in carbide
content in the steel, resulting in a decrease in ductility, and
hence a decrease in press formability. A more serious problem is
that a C content exceeding 0.20% leads to remarkable deterioration
of spot weldability and arc weldability. For these reasons, in the
present invention, the C content is limited to not more than 0.20%.
From the viewpoint of formability, the C content is preferably not
more than 0.18%.
[0261] Si: not more than 2.0%
[0262] Si is a useful strengthening element, which improves
strength of a steel sheet without a marked decrease in ductility of
the steel sheet, and is necessary for obtaining retained austenite.
These effects are particularly remarkable at an Si content of not
less than 0.1% and therefore, the Si content is preferably not less
than 0.1%. An Si content exceeding 2.0%, however, leads to
deterioration of press formability and degrades the platability.
Therefore, the Si content is limited to not more than 2.0%.
[0263] Mn: not more than 3.0%
[0264] Mn is a useful element, which strengthens the steel and
prevents hot cracking caused by S, and is therefore contained in an
amount according to S content. These effects are particularly
remarkable at an Mn content of not less than 0.5%. However, an Mn
content exceeding 3.0% results in deterioration of press
formability and weldability. The Mn content is, therefore, limited
to not more than 3.0%. More preferably, the Mn content is not less
than 1.0%.
[0265] P: not more than 0.10%
[0266] P strengthens the steel. In the present invention, P is
preferably contained in an amount of not less than 0.005% for
securing the strength. However, an excess content of P exceeding
0.10% causes deterioration of press formability. For this reason,
in the present invention, a P content is limited to not more than
0.10%. When more enhanced press formability is required, the P
content is preferably not more than 0.08%.
[0267] S.: not more than 0.02%
[0268] S is an element, which is present as inclusions in a steel
sheet and causes deterioration of ductility, formability, and
particularly stretch flanging formability of the steel sheet, and
it should be the lowest possible. An S content reduced to not more
than 0.02% does not exert much adverse effect and therefore, the. S
content is limited to not more than 0.02% in the present invention.
When excellent stretch flanging formability is required, the S
content is preferably not more than 0.010%.
[0269] Al: not more than 0.10%
[0270] Al is a deoxidizing element of steel, and is useful for
improving cleanliness of steel. In addition, Al is effective for
the formation of the retained austenite. In the present invention,
the Al content is preferably not less than 0.01%. An excess Al
content exceeding 0.30%, however, cannot give a further enhanced
effect because of saturation of the effect, and causes
deterioration of press formability. The Al content is, therefore,
limited to not more than 0.30%. The present invention also include
a steel making process using other deoxidizers, for example, Ti or
Si, and steel sheets produced by such deoxidation methods are also
included in the scope of the present invention. In this case,
addition of Ca or REM to molten steel does not impair the features
of the steel sheet of the present invention at all. Of course,
steel sheets containing Ca or REM are included within the scope of
the present invention.
[0271] N: not more than 0.02%
[0272] N is an element, which increases strength of a steel sheet
through solid solution strengthening or strain age hardening, and
is preferably contained in an amount of not less than 0.001%. An N
content exceeding 0.02% causes an increase in the nitride content
in the steel sheet, which causes serious deterioration of ductility
and of press formability. The N content is, therefore, limited to
not more than 0.02%. When further improvement of press formability
is required, the N content is preferably not more than 0.01%.
[0273] Cu: 0.5 to 3.0%
[0274] Cu is an element, which remarkably increases strain age
hardening of a steel sheet (increase in strength after
pre-deformation/heat treatment), and is the most important element
in the present invention. With a Cu content of less than 0.5%, an
increase in tensile strength .DELTA.TS of not less than 80 MPa
cannot be obtained by changing the pre-deformation/heat treatment
conditions. In the present invention, therefore, Cu should be
contained in an amount of not less than 0.5%. With a Cu content
exceeding 3.0%, however, the effect is saturated, leading to
unfavorable economic effects. Furthermore, deterioration of press
formability occurs, and the surface quality of the steel sheet is
degraded. The Cu content is, therefore, limited within the range of
0.5 to 3.0%. In order to simultaneously achieve a higher .DELTA.TS
and excellent press formability, the Cu content is preferably
within the range of 1.0 to 2.5%.
[0275] In the present invention, it is preferable that the
composition containing Cu further contain, in weight percent, at
least one of the following Groups A to C:
[0276] Group A: Ni: not more than 2.0%;
[0277] Group B: at least one of Cr and Mo: not more than 2.0% in
total; and
[0278] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total.
[0279] Group A: Ni: not more than 2.0%
[0280] Group A: Ni is an element effective for preventing surface
defects produced by Cu contained in the steel sheet, and can be
contained as required. The Ni content depends on the Cu content,
and is preferably about a half the Cu content, more specifically,
within the range of about 30 to about 80% of the Cu content. An Ni
content exceeding 2.0% cannot give further enhancement in the
effect because of saturation of the effect, leading to economic
disadvantages, and causes deterioration of press formability. For
these reasons, the Ni content is preferably limited to not more
than 2.0%.
[0281] Group B: at least one of Cr and Mo: not more than 2.0% in
total
[0282] Group B: Both Cr and Mo strengthen the steel sheet, like Mn,
and can be contained as required. However, if at least one of Cr
and Mo are contained in an amount exceeding 2.0% in total, press
formability is impaired. The total content of Cr and Mo is
preferably limited to not more than 2.0%. From the viewpoint of
press formability, a Cr content is preferably not less than 0.1%,
and an Mo content, is preferably not less than 0.1%.
[0283] Group C: at least one of Nb, Ti, and V: not more than 0.2%
in total
[0284] Group C: Nb, Ti, and V are carbide-forming elements and
increase strength by fine dispersion of carbides, and can be
selected and contained as required. However, if the total content
of at least one of Nb, Ti, and V exceeds 0.2%, press formability is
impaired. Thus, the total content of Nb, Ti and V is preferably
limited to not more than 0.2%. The above-mentioned effect can be
achieved at an Nb content of not less than 0.01%, at a Ti content
of not less than 0.01%, and at a V content of not less than
0.01%.
[0285] In the present invention, in place of Cu, at least one
selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to
2.0%, Cr, and W: 0.05 to 2.0% may be contained in an amount of not
more than 2.0% in total.
[0286] At least one selected from the group consisting of Mo: 0.05
to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not
more than 2.0% in total
[0287] In the present invention, all of Mo, Cr, and W, as well as
Cu, are the most important elements, which remarkably increase
strain age hardening (increase in strength after
pre-deformation/heat treatment) of the steel sheet. When a steel
sheet containing at least one of Mo, Cr, and W, and having a
composite structure comprising a primary phase of a ferrite phase
and a tempered martensite phase and a secondary phase containing
retained austenite in a volume ratio of not less than 1% is
subjected to prestrain (pre-deformation) of not less than 5% and a
low-temperature heat treatment (heat treatment), the retained
austenite is transformed into martensite by strain-induced
transformation. Then, the formation of fine carbide precipitation
is induced by the strain at a low temperature occurs in the
martensite, resulting in an increase in tensile strength .DELTA.TS
of not less than 80 MPa. With a content of each of these elements
of less than 0.05%, changing the steel sheet structure and
pre-deformation/heat treatment conditions does not give an increase
in tensile strength .DELTA.TS of not less than 80 MPa. Therefore,
in the present invention, each of Mo, Cr, and W is preferably
contained in an amount of not less than 0.05%. If the content of
each of Mo, Cr, and W each exceeds 2.0%, a further enhanced effect
corresponding to the content cannot be expected as a result of
saturation of the effect, leading to economic disadvantages, and
this results in deterioration of press formability. For these
reasons, the content of each of Mo, Cr, and W is preferably limited
within the range of 0.05 to 2.0%, and the total content thereof is
preferably limited to not more than 2.0%.
[0288] The above-mentioned composition containing at least one of
Mo, Cr, and W preferably further contains at least one of Nb, Ti,
and V in an amount of not more than 2.0% in total.
[0289] At least one of Nb, Ti, and V, in a total amount of not more
than 2.0%
[0290] Nb, Ti, and V are carbide-forming elements and can be
selected and contained as required, when at least one of Mo, Cr,
and W is added. However, a total content of Nb, Ti, and V exceeding
2.0% causes deterioration of press formability. Thus, the total
content of Nb, Ti, and V is preferably limited to not more than
2.0%. At least one of Mo, Cr, and W are added, at least one of Nb,
Ti, and V are added, and the structure is transformed into a
composite structure of a primary phase comprising a ferrite phase
and a tempered martensite phase and a secondary phase containing
retained austenite. This forms fine composite carbides in
martensite which was formed by strain-induced transformation during
the pre-deformation/heat treatment, and strain-induced fine
precipitation at a low temperature occurs, resulting in an increase
in tensile strength .DELTA.TS of not less than 80 MPa. In order to
obtain this effect, Nb, Ti, and V is preferably contained in an
amount of not less than 0.01% for Nb, in an amount of not less than
0.01% for Ti and in an amount of not less than 0.01% for V, and at
least one of Nb, Ti, and V can be selected and contained as
required.
[0291] Although no particular restriction is imposed, apart from
the above-mentioned constituents, the composition, may contain B:
not more than 0.1%, Ca: not more than 0.1%, Zn: not more than 0.1%,
and REM: not more than 0.1% without any problem.
[0292] The balance of the composition of the steel is Fe and
incidental impurities. Allowable incidental impurities include Sb:
not more than 0.01%, Sn: not more than 0.1%, Zn: not more than
0.01%, and Co: not more than 0.1%.
[0293] The method for manufacturing the hot-dip galvanized steel
sheet of the present invention will now be described.
[0294] The hot-dip galvanized steel sheet is preferably
manufactured through a primary heat treatment step of heating a
steel sheet having the above-mentioned composition to a temperature
of not less than the A.sub.C1 transformation point and rapidly
cooling the steel sheet, a secondary heat treatment step of heating
the steel sheet to a temperature of ferrite/austenite dual phase
within the range of A.sub.C1 transformation point to A.sub.C3
transformation point on a continuous hot-dip galvanizing line, and
a hot-dip galvanizing step of forming a hot-dip galvanizing layer
on each surface of the steel sheet.
[0295] A hot-rolled steel sheet or a cold-rolled steel sheet may
preferably be used in this process. A preferable manufacturing
method of the steel sheet used will now be described, although the
method is not limited thereto in the present invention.
[0296] A suitable method for manufacturing the hot-rolled steel
sheet used as a galvanizing substrate will be described.
[0297] A material (steel slab) used is preferably manufactured by a
continuous casting process to prevent macro-segregation of the
constituents, but it may be manufactured by an ingot casting
process or a thin-slab casting process. A conventional process
employed in this embodiment includes the steps of manufacturing a
steel slab, cooling the steel slab to room temperature, and
reheating the slab. Alternatively, an energy-saving process is
applicable with no problem. As the energy-saving process, for
example, a direct-hot charge rolling process of charging the hot
steel slab into a reheating furnace without cooling the same, and a
direct rolling process of immediately rolling after a short
temperature holding are applicable.
[0298] The material (steel slab) is first heated, and subjected to
a hot rolling step to form a hot-rolled steel sheet. Known hot
rolling conditions may be employed without problem as long as a
hot-rolled steel sheet having a desired thickness is formed.
Preferable conditions for hot rolling are as follows:
[0299] Slab reheating temperature: not less than 900.degree. C.
[0300] In the case of a steel slab containing Cu, the slab heating
temperature is preferably the lowest possible to prevent surface
defects caused by Cu. However, a heating temperature of less than
900.degree. C. causes an increase in the rolling load, thus
increasing the risk of occurrence of a trouble during the hot
rolling. Considering the increase in scale loss caused by
accelerated oxidation, the slab heating temperature is preferably
not more than 1,300.degree. C. From the viewpoint of decreasing the
slab heating temperature and preventing occurrence of troubles
during hot rolling, use of a so-called sheet bar heater, which
heats a sheet bar, is effective.
[0301] Finish rolling end temperature: not less than 700.degree.
C.
[0302] At a finish rolling end temperature FDT of not less than
700.degree. C., it is possible to obtain a uniform hot-rolled
mother sheet structure which can give an excellent formability
after cold rolling and recrystallization annealing. A finish
rolling end temperature FDT of less than 700.degree. C. leads to a
non-uniform structure of the hot-rolled mother sheet and a higher
rolling load during hot rolling, thus increasing the risk of
occurrence of troubles during hot rolling. Thus, the FDT for the
hot rolling step is preferably not less than 700.degree. C.
[0303] Coiling temperature: not more than 800.degree. C.
[0304] The coiling temperature CT is preferably not more than
800.degree. C., and more preferably not less than 200.degree. C.
The CT exceeding 800.degree. C. tends to cause a decrease in yield
as a result of an increased scale loss. With a CT of less than
200.degree. C., the steel sheet shape is seriously impaired, and
there is an increasing risk of occurrence of inconveniences in
practical use.
[0305] The hot-rolled steel sheet suitably applicable in the
present invention is preferably prepared by heating the slab to not
less than 900.degree. C., hot-rolling the heated slab at a finish
rolling end temperature of not less than 700.degree. C., and
coiling the hot-rolled sheet at a coiling temperature of not less
than 800.degree. C., and preferably not less than 200.degree.
C.
[0306] In the above-mentioned hot rolling step, all or part of
finish rolling may be lubrication rolling, which reduces the
rolling load during the hot rolling. The lubrication rolling is
effective also from the viewpoint of achieving a uniform steel
sheet shape and a uniform material quality. The frictional
coefficient on the lubrication rolling is preferably within the
range of 0.25 to 0.10. It is desirable to connect neighboring sheet
bars to each other to perform a continuous finish rolling process.
Application of the continuous rolling process is desirable also
from the viewpoint of operational stability of hot rolling.
[0307] The hot-rolled sheet with scales may be annealed to form an
internal oxide layer at the surface of the steel sheet. The
internal oxide layer, which prevents concentration of Si, Mn, and P
at the surface, improves hot-dip galvanizing ability.
[0308] The hot-rolled sheet manufactured by the above-mentioned
method may be used as an original sheet for plating. Alternatively,
the hot-rolled sheet may be cold-rolled to form a cold-rolled sheet
used as an original sheet for plating.
[0309] In the cold rolling step, any cold rolling condition may be
used without particular restriction so far as such a condition
permits production of cold-rolled steel sheets with desired
dimensions and shapes. The reduction in cold rolling is preferably
not less than 40%. A reduction of less than 40% inhibits uniform
recrystallization during the subsequent primary heat treatment.
[0310] In the present invention, the above-mentioned steel sheet
(hot-rolled sheet or cold-rolled sheet) is subjected to a primary
heat treatment step including heating to a temperature of not less
than the A.sub.C1 transformation point and rapid cooling.
[0311] Heating in the primary heat treatment, the steel sheet is
preferably held at a temperature of not less than A.sub.C1
transformation point, more preferably not less than (A.sub.C3
transformation point -50.degree. C.), and most preferably not less
than A.sub.C3 transformation point. After heating, the steel sheet
is preferably rapidly cooled to a temperature of not more than the
Ms point at a cooling rate of not less than 10.degree. C./second.
During the primary heat treatment step, lath martensite is produced
in the steel sheet. In the present invention, the most important
point is to form lath martensite during the primary heat treatment
step. Unless the lath martensite is formed in the steel sheet, it
is difficult to form a secondary phase containing retained
austenite in the subsequent steps.
[0312] When a hot-rolled steel sheet, subjected to final hot
rolling at a temperature of not less than (Ar.sub.3 transformation
point -50.degree. C.), is used as an original sheet for plating,
the primary heat treatment step can be substituted the steel sheet
for rapidly cooling to a temperature of not less than Ms point at a
cooling rate of not less than 10.degree. C./second during cooling
after the final hot rolling.
[0313] Then, the steel sheet containing lath martensite formed
during the above-described primary heat treatment is subjected to a
secondary heat treatment step for heating to and holding at a
temperature in the range of A.sub.C1 transformation point to
A.sub.C3 transformation point on a continuous hot-dip galvanizing
line. During the secondary heat treatment step, the lath martensite
formed during the primary heat treatment step is changed into
tempered martensite, and a part of the structure is transformed
into austenite for formation of retained austenite.
[0314] A heating and holding temperature of less than the A.sub.C1
transformation point in the secondary heat treatment step cannot
form retained austenite. A heating and holding temperature
exceeding the A.sub.C3 transformation point causes retransformation
of the entire structure of the steel sheet to austenite, whereby
the tempered martensite disappears. For these reasons, the heating
and holding temperature in the secondary heat treatment is within
the range of the A.sub.C1 transformation point to the A.sub.C3
transformation point.
[0315] Then, the steel sheet heated to and held at a temperature in
the range of the A.sub.C1 transformation point to the A.sub.C3
transformation point in the second heat treatment step is
preferably cooled to a temperature of not more than 500.degree. C.
at a cooling rate of 5.degree. C./second or more, from the
viewpoint of forming retained austenite. This can achieve a
composite structure of a primary phase containing a ferrite phase
and a tempered martensite phase and a secondary phase containing
retained austenite in the steel sheet.
[0316] The steel sheet after the secondary heat treatment is
subsequently subjected to a hot-dip galvanizing treatment step on a
continuous hot-dip galvanizing line.
[0317] The hot-dip galvanizing treatment may be carried out under
treatment conditions (galvanizing bath temperature: 450 to
500.degree. C.) used in a usual continuous hot-dip galvanizing line
without a particular restriction. Because galvanizing at an
excessively high temperature leads to a poor platability,
galvanizing is preferably conducted at a temperature of not more
than 500.degree. C. Galvanizing at a temperature of less than
450.degree. C. causes deterioration of platability. From the
viewpoint of forming martensite, the cooling rate from the hot-dip
galvanizing temperature to 300.degree. C. is preferably not less
than 5.degree. C./second.
[0318] For the purpose of adjusting the galvanizing weight as
required after galvanizing, wiping may be performed.
[0319] After the hot-dip galvanizing treatment, an alloying
treatment of a galvanizing layer may be applied. The alloying
treatment is preferably carried out by reheating the plated sheet
to a temperature in the range of 450 to 500.degree. C. after the
hot-dip galvanizing treatment. At an alloying treatment temperature
of less than 450.degree. C., alloying is decelerated, resulting in
low productivity. On the other hand, an alloying treatment
temperature exceeding 550.degree. C. causes deterioration of
platability, makes it difficult to secure a required amount of
retained austenite, and decrease ductility of the steel sheet.
[0320] After the alloying treatment, the sheet is preferably cooled
to 300.degree. C. at a cooling rate of not less than 5.degree.
C./second. An extremely low cooling rate after the alloying
treatment makes it difficult to form a required amount of retained
austenite.
[0321] In the present invention, pickling treatment for removing a
concentrated surface layer of the constituents formed on the
surface of the steel sheet during the primary heat treatment step
is preferably performed between the primary heat treatment step and
the hot-dip galvanizing step, for the improvement in platability.
By the primary heat treatment, P and oxides of Si, Mn, Cr, etc. are
concentrated on the steel surface to form a concentrated surface
layer. It is favorable for improving platability to remove this
concentrated surface layer through pickling and to conduct
annealing in a reducing atmosphere subsequently on the continuous
hot-dip galvanizing line.
[0322] After the hot-dip galvanizing or the alloying treatment
step, a temper rolling step with a reduction of not more than 10%
may be applied for adjustments such as shape correction and surface
roughness adjustment.
[0323] To the steel sheet of the present invention, any special
treatment may be applied after the hot-dip galvanizing, to improve
chemical treatment ability, weldability, press formability, and
corrosion resistance.
EXAMPLES
Example 1
[0324] Molten steels having the compositions shown in Table 1 were
made in a converter and cast into steel slabs by a continuous
casting process. Each of these steel slabs was reheated, and
hot-rolled under conditions shown in Table 2 into a hot-rolled
steel strip (hot-rolled sheet) having a thickness of 2.0 mm. The
hot-rolled sheet was temper-rolled at a reduction of 1.0%.
1TABLE 1 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N A 0.09 1.45
1.05 0.01 0.003 0.034 0.002 B 0.12 1.50 1.20 0.01 0.002 0.030 0.002
C 0.10 1.48 1.35 0.01 0.002 0.028 0.002 D 0.15 1.53 1.45 0.01 0.003
0.033 0.002 E 0.12 1.48 1.55 0.01 0.005 0.032 0.002 F 0.11 1.50
1.08 0.01 0.004 0.032 0.002 G 0.13 1.52 1.22 0.01 0.004 0.030 0.002
H 0.12 1.42 1.22 0.01 0.003 0.033 0.002 I 0.11 1.52 1.52 0.01 0.003
0.031 0.002 J 0.13 1.43 1.48 0.01 0.003 0.028 0.002 K 0.15 1.58
1.05 0.01 0.003 0.030 0.002 L 0.14 1.60 1.21 0.01 0.003 0.028 0.002
STEEL COMPOSITION NO. Cu Ni Cr, Mo, Nb, Ti, V A 1.52 -- -- -- B
1.43 0.65 Mo: 0.32 -- C 1.25 0.52 Cr: 0.53 -- D 1.33 0.44 -- Nb:
0.01, Ti: 0.01, V: 0.01 E 0.15 -- -- -- F 0.68 -- -- -- G 0.98 --
-- -- H 1.55 0.62 -- -- I 1.49 -- Cr: 0.15, -- Mo: 0.12 J 1.43 --
Mo: 0.21 -- K 1.52 -- -- Nb: 0.01 L 1.48 -- Cr: 0.11 Ti: 0.01
[0325]
2 TABLE 2 HOT ROLLING - COOLING AFTER ROLLING SLAB FINISH FORCED
COOLING ISOTHERMAL STEEL REHEATING ROLLING TIME BEFORE COOLING
HOLDING SHEET STEEL TEMP. END TEMP. START RATE STOP TEMP. HOLDING
NO. NO. SRT .degree. C. FDT .degree. C. COOLING S .degree. C./s
TEMP. .degree. C. .degree. C. TIME S 1 A 1250 850 0.5 100 710 710 5
2 B 1250 850 0.5 80 690 690 5 3 1250 850 0.3 30 700 -- -- 4 1250
850 0.5 30 680 -- -- 5 C 1250 850 0.1 60 700 700 5 6 D 1250 850 0.5
80 680 680 5 7 E 1250 850 0.5 70 710 710 5 8 F 1250 850 0.5 60 700
700 5 9 G 1250 850 0.5 80 690 690 5 10 H 1250 850 0.5 60 680 680 5
11 I 1250 850 0.1 60 690 -- -- 12 J 1250 850 0.1 80 700 -- -- 13 K
1250 850 0.1 80 680 680 5 14 L 1250 850 0.3 60 690 690 5 15 H 1250
750 0.5 50 620 620 5 16 1250 850 3.0 50 680 680 12 17 1250 850 0.5
30 680 680 5 18 1250 850 0.5 60 600 600 5 19 1250 850 0.5 60 700 --
-- HOT ROLLING - COOLING AFTER ROLLING STEEL SLOW COOLING TREATMENT
COOLING COILING SHEET INITIAL COOLING STOP RATE BEFORE TEMP. NO.
TEMP. .degree. C. RATE .degree. C./s TEMP. .degree. C. COILING
.degree. C/s CT .degree. C. 1 -- -- -- 80 450 2 -- -- -- 60 450 3
700 10 650 30 600 4 680 10 650 20 450 5 -- -- -- 60 450 6 -- -- --
80 450 7 -- -- -- 80 450 8 -- -- -- 70 450 9 -- -- -- 80 450 10 --
-- -- 60 450 11 690 10 650 60 450 12 700 10 650 60 450 13 680 10
640 80 450 14 690 10 650 60 450 15 620 10 580 60 450 16 -- -- -- 70
450 17 680 10 650 60 450 18 -- -- -- 70 450 19 -- -- -- 70 450
[0326] For the resulting hot-rolled steel strip (hot-rolled steel
sheet), the microstructure, tensile properties, strain age
hardenability, and hole expanding property were determined. Press
formability was evaluated in terms of elongation El (ductility),
TS.times.El balance and hole expanding ratio .lambda.. Test methods
were as follows.
[0327] (1) Microstructure
[0328] A test piece was sampled from each of the resultant
hot-rolled sheets, and the microstructure of the cross-section
(section C) perpendicular to the rolling direction of the steel
sheet was observed with an optical microscope and a scanning
electron microscope. The volume ratios of the ferrite phase, the
bainite phase, and the martensite phase in the steel sheet were
determined with an image analyzer using a photograph of the
cross-sectional structure at a magnification of 1,000. The volume
ratios of the retained austenite phase were determined by polishing
the steel sheet to the central plane in the thickness direction,
and by measuring diffraction X-ray intensities at the central
plane. Mo K.alpha.-rays were used as incident X-rays, the ratios of
the diffraction X-ray strengths of the planes {200}, {220} and
{311} of the retained austenite phase to the diffraction X-ray
strengths of the planes {110}, {200} and {211} of the ferrite
phase, respectively, were determined, and the volume ratio of the
retained austenite was determined from the average of these
ratios.
[0329] (2) Tensile Properties
[0330] JIS No. 5 tensile test pieces were sampled from the
resultant hot-rolled sheets, and a tensile test was carried out in
accordance with JIS Z 2241 to determine the yield strength YS, the
tensile strength TS, and the elongation El.
[0331] (3) Strain Age Hardenability
[0332] JIS No. 5 test pieces were sampled in the rolling direction
from the resultant hot-rolled steel sheets. A plastic deformation
of 5% was applied as a pre-deformation (tensile prestrain). After a
heat treatment at 250.degree. C. for 20 minutes, a tensile test was
carried out to determine tensile properties (yield stress YS.sub.TH
and tensile strength TS.sub.HT) and to calculate
.DELTA.YS=YS.sub.TH-YS, and .DELTA.TS=TS.sub.HT-TS, wherein
YS.sub.TH and TS.sub.HT were yield stress and tensile strength
after the pre-deformation/heat treatment, and YS and TS were yield
stress and tensile strength of the hot-rolled steel sheets.
[0333] (4) Hole Expanding Property
[0334] A hole was formed by punching a test piece sampled from the
resultant hot-rolled sheet in accordance with Japan Iron and Steel
Federation Standard (JFS T 1001-1996) with a punch having a
diameter of 10 mm. Then, the hole was expanded with a conical punch
having a vertical angle of 600 so that burrs were produced on the
outside until cracks passing through the thickness form, thereby
determining the hole expanding ratio .lambda.. The hole expanding
ratio .lambda. was calculated by the formula:
.lambda.(%)={(d-d.sub.0)/d.sub.0}.times.100, where d.sub.0 is
initial hole diameter (punch diameter), and d is inner hole
diameter upon occurrence of cracks.
[0335] The results are shown in Table 3.
3 TABLE 3 MICROSTRUCTURE PRIMARY SECONDARY PHASE HOT-ROLLED SHEET
PROPERTIES STEEL PHASE OTHER TENSILE PROPERTIES SHEET STEEL F
VOLUME A VOLUME PHASES VOLUME YS TS El TS .times. El NO. NO. RATIO
% RATIO % KIND* RATIO % (MPa) (MPa) % MPa % 1 A 75 8 B, M 25 470
620 34 21080 2 B 80 11 B, M 20 490 650 33 21450 3 75 -- P 25 660
720 15 10800 4 76 -- P, B 24 600 660 16 10560 5 C 78 9 B, M 22 490
650 33 21450 6 D 75 9 B, M 25 500 660 32 21120 7 E 80 8 B, M 20 410
540 39 21060 8 F 81 10 B, M 19 470 620 34 21080 9 G 80 9 B, M 20
460 610 35 21350 10 H 80 9 B, M 20 490 650 33 21450 11 I 81 8 B, M
19 470 620 34 21080 12 J 78 10 B, M 22 500 660 32 21120 13 K 80 8
B, M 20 470 620 34 21080 14 L 75 10 B, M 25 500 660 32 21120 15 H
80 -- P, B 20 600 660 16 10560 16 80 -- P 20 590 650 15 9750 17 80
-- P, B 20 610 670 14 9380 18 80 -- P, 20 580 640 17 10880 19 78 --
P, B 22 590 650 15 9750 PROPERTIES AFTER PRE- STRAIN AGE HOLE
DEFORMATION - HARDENING EXPANSION STEEL HEAT TREATMENT PROPERTIES
HOLE SHEET YS.sub.HT TS.sub.HT .DELTA.YS .DELTA.TS EXPANDING NO.
MPa MPa MPa MPa RATIO .lambda. % REMARKS 1 715 790 245 170 140
EXAMPLE 2 750 830 260 180 135 EXAMPLE 3 730 760 70 40 70 COMP. EX.
4 660 695 60 35 60 COMP. EX. 5 730 810 240 160 145 EXAMPLE 6 745
825 245 165 140 EXAMPLE 7 715 550 305 10 60 COMP. EX. 8 675 750 205
130 140 EXAMPLE 9 690 765 230 155 135 EXAMPLE 10 750 830 260 180
135 EXAMPLE 11 675 750 205 130 140 EXAMPLE 12 745 825 245 165 140
EXAMPLE 13 715 790 245 170 140 EXAMPLE 14 745 825 245 165 140
EXAMPLE 15 660 695 60 35 60 COMP. EX. 16 660 690 70 40 70 COMP. EX.
17 670 705 60 35 70 COMP. EX. 18 650 675 70 35 60 COMP. EX. 19 650
690 60 40 70 COMP. EX. *F: FERRITE, A: AUSTENITE, M: MARTENSITE, P:
PEARLITE, B: BAINITE
[0336] All Examples according to the present invention show a high
elongation El, a high strength/ductility balance (TS.times.El), and
a high hole expanding ratio .lambda., suggesting excellent stretch
flanging formability. In addition, all Examples according to the
present invention show a very large .DELTA.TS, suggesting that
these samples had excellent strain age hardenability. Comparative
Examples outside the scope of the present invention, in contrast,
suggest that the samples have a low elongation El, a small hole
expanding ratio .lambda., a low .DELTA.TS, and decreased press
formability and strain age hardenability.
Example 2
[0337] Molten steels having the compositions shown in Table 4 were
made in a converter and cast into steel slabs by a continuous
casting process. Each of these steel slabs were reheated, and
hot-rolled under conditions shown in Table 5 into a hot-rolled
steel strip (hot-rolled sheet) having a thickness of 2.0 mm. The
hot-rolled steel strip was temper-rolled at a reduction of
1.0%.
4TABLE 1 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N A 0.09 1.45
1.05 0.01 0.003 0.034 0.002 B 0.12 1.50 1.20 0.01 0.002 0.030 0.002
C 0.10 1.48 1.35 0.01 0.002 0.028 0.002 D 0.15 1.53 1.45 0.01 0.003
0.033 0.002 E 0.12 1.48 1.55 0.01 0.005 0.032 0.002 F 0.11 1.50
1.08 0.01 0.004 0.032 0.002 G 0.13 1.52 1.22 0.01 0.004 0.030 0.002
H 0.12 1.42 1.22 0.01 0.003 0.033 0.002 I 0.11 1.52 1.52 0.01 0.003
0.031 0.002 J 0.13 1.43 1.48 0.01 0.003 0.028 0.002 K 0.15 1.58
1.05 0.01 0.003 0.030 0.002 L 0.14 1.60 1.21 0.01 0.003 0.028 0.002
A 0.09 1.45 1.05 0.01 0.003 0.034 0.002 B 0.12 1.50 1.20 0.01 0.002
0.030 0.002 C 0.10 1.48 1.35 0.01 0.002 0.028 0.002 D 0.15 1.53
1.45 0.01 0.003 0.033 0.002 E 0.12 1.48 1.55 0.01 0.005 0.032 0.002
F 0.11 1.50 1.08 0.01 0.004 0.032 0.002 G 0.13 1.52 1.22 0.01 0.004
0.030 0.002 H 0.12 1.42 1.22 0.01 0.003 0.033 0.002 I 0.11 1.52
1.52 0.01 0.003 0.031 0.002 J 0.13 1.43 1.48 0.01 0.003 0.028 0.002
K 0.15 1.58 1.05 0.01 0.003 0.030 0.002 L 0.14 1.60 1.21 0.01 0.003
0.028 0.002 STEEL COMPOSITION NO. Cu Ni Cr, Mo, Nb, Ti, V A 1.52 --
-- -- B 1.43 0.65 Mo: 0.32 -- C 1.25 0.52 Cr: 0.53 -- D 1.33 0.44
-- Nb: 0.01, Ti: 0.01, V: 0.01 E 0.15 -- -- -- F 0.68 -- -- -- G
0.98 -- -- -- H 1.55 0.62 -- -- I 1.49 -- Cr: 0.15, -- Mo: 0.12 J
1.43 -- Mo: 0.21 -- K 1.52 -- -- Nb: 0.01 L 1.48 -- Cr: 0.11 Ti:
0.01 A 1.52 -- -- -- B 1.43 0.65 Mo: 0.32 -- C 1.25 0.52 Cr: 0.53
-- D 1.33 0.44 -- Nb: 0.01, Ti: 0.01, V: 0.01 E 0.15 -- -- -- F
0.68 -- -- -- G 0.98 -- -- -- H 1.55 0.62 -- -- I 1.49 -- Cr: 0.15,
-- Mo: 0.12 J 1.43 -- Mo: 0.21 -- K 1.52 -- -- Nb: 0.01 L 1.48 --
Cr: 0.11 Ti: 0.01
[0338] For the resultant hot-rolled steel strip (hot-rolled steel
sheet), the microstructure, the tensile properties, the strain age
hardenability, and the hole expanding ratio were determined as in
Example 1. Press formability was evaluated in terms of elongation
El (ductility), TS.times.El balance and the hole expanding ratio
.lambda..
[0339] The results obtained are shown in Table 6.
5 TABLE 6 MICROSTRUCTURE PRIMARY SECONDARY PHASE HOT-ROLLED SHEET
PROPERTIES STEEL PHASE OTHER TENSILE PROPERTIES SHEET STEEL F
VOLUME A VOLUME PHASES VOLUME YS TS El TS .times. El NO. NO. RATIO
% RATIO % KIND* RATIO % (MPa) (MPa) % MPa % 2-1 2A 76 8 B, M 24 460
610 35 21350 2-2 2B 79 9 B, M 21 480 640 33 21120 2-3 76 -- P 24
650 710 15 10650 2-4 75 -- P, B 25 590 650 14 9100 2-5 2C 76 9 B, M
24 480 630 34 21420 2-6 2D 78 8 B, M 22 490 650 33 21450 2-7 2E 80
7 B, M 20 390 510 42 21420 2-8 2F 81 9 B, M 19 450 590 36 21240 2-9
2G 79 10 B, M 21 450 600 36 21600 2-10 2H 78 10 B, M 22 480 630 34
21420 2-11 2I 80 8 B, M 20 460 610 35 21350 2-12 2J 79 9 B, M 21
450 590 36 21240 2-13 2K 80 9 B, M 20 460 600 35 21000 2-14 2L 81 8
B, M 19 470 620 34 21080 PROPERTIES AFTER PRE- STRAIN AGE HOLE
DEFORMATION - HARDENING EXPANSION STEEL HEAT TREATMENT PROPERTIES
HOLE SHEET YS.sub.HT TS.sub.HT .DELTA.YS .DELTA.TS EXPANDING NO.
MPa MPa MPa MPa RATIO .lambda. % REMARKS 2-1 695 760 235 150 135
EXAMPLE 2-2 730 800 250 160 140 EXAMPLE 2-3 700 730 50 20 70 COMP.
EX. 2-4 635 665 45 15 65 COMP. EX. 2-5 715 785 235 155 140 EXAMPLE
2-6 725 810 235 160 135 EXAMPLE 2-7 620 670 230 160 130 EXAMPLE 2-8
660 730 210 140 135 EXAMPLE 2-9 570 630 120 30 65 COMP. EX. 2-10
715 785 235 155 130 EXAMPLE 2-11 695 760 235 150 135 EXAMPLE 2-12
660 730 210 140 130 EXAMPLE 2-13 670 750 200 150 140 EXAMPLE 2-14
670 780 200 160 135 EXAMPLE *F: FERRITE, A: AUSTENITE, M:
MARTENSITE, P: PEARLITE, B: BAINITE
[0340] All Examples according to the present invention showed a
high elongation El, a high strength-ductility balance (TS.times.El)
having excellent press formatility, and further showed a very large
.DELTA.TS, suggesting that these samples had excellent strain age
hardenability. Comparative Examples outside the scope of the
present invention, in contrast, suggest that the samples had a low
elongation El, a low .DELTA.TS, and decreased press formability and
strain age hardenability.
Example 3
[0341] Molten steels having the composition shown in Table 7 were
made in a converter and cast into steel slabs by a continuous
casting process. Then, each of these steel slabs was reheated to
1,250.degree. C., and hot-rolled in a hot rolling step of hot
rolling at a finish rolling end temperature of 900.degree. C. and a
coiling temperature of 600.degree. C. into a hot-rolled steel strip
(hot-rolled sheet) having a thickness of 4.0 mm. Then, the
hot-rolled steel strip (hot-rolled sheet) was subjected to a cold
rolling step of pickling and cold-rolling into cold rolled steel
strip (cold-rolled sheet) having a thickness of 1.2 mm. Thereafter,
the cold-rolled steel strip (cold-rolled sheet) was subjected to
recrystallization annealing step comprising heating and soaking
treatment and a subsequent retaining treatment under the conditions
shown in Table 8 on the continuous annealing line to obtain
cold-rolled annealed sheet. The resultant steel strip (cold-rolled
annealed sheet) was further temper-rolled at an reduction of
0.8%.
6TABLE 7 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N Cu Ni 3A
0.10 1.20 1.42 0.01 0.003 0.032 0.002 1.51 -- 3B 0.11 1.10 1.51
0.01 0.002 0.033 0.002 1.45 0.63 3C 0.11 1.32 1.33 0.01 0.004 0.025
0.002 1.20 0.52 3D 0.10 1.06 1.48 0.01 0.003 0.022 0.002 1.39 0.43
3E 0.09 1.25 1.36 0.01 0.004 0.029 0.002 0.22 -- 3F 0.10 1.08 1.45
0.01 0.001 0.030 0.002 0.75 -- 3G 0.11 1.15 1.52 0.01 0.002 0.033
0.002 0.96 -- 3H 0.10 1.10 1.55 0.01 0.002 0.025 0.002 1.22 0.66 3I
0.11 1.09 1.48 0.01 0.001 0.033 0.002 1.36 -- 3J 0.11 1.12 1.62
0.01 0.002 0.029 0.001 1.42 -- 3K 0.10 1.25 1.39 0.01 0.002 0.032
0.002 1.38 -- 3L 0.09 1.10 1.45 0.01 0.003 0.025 0.002 1.29 -- 3M
0.10 1.35 1.50 0.01 0.002 0.030 0.002 1.44 -- 3N 0.11 1.26 1.46
0.01 0.001 0.028 0.001 1.33 0.52 TRANSFORMATION STEEL COMPOSITION
POINT (.degree. C.) NO. Cr Mo Nb Ti V Ac1 Ac3 3A -- -- -- -- -- 725
875 3B -- 0.11 -- -- -- 715 875 3C 0.12 -- -- -- -- 725 880 3D --
-- 0.01 0.01 0.01 720 870 3E -- -- -- -- -- 730 860 3F -- -- -- --
-- 720 880 3G -- -- -- -- -- 725 875 3H -- -- -- -- -- 730 875 3I
-- 0.10 -- -- -- 725 860 3J 0.10 -- -- -- -- 730 880 3K -- -- 0.01
-- -- 720 870 3L -- -- -- 0.01 -- 725 865 3M -- -- -- -- 0.01 730
875 3N 0.12 0.11 0.01 0.01 0.01 725 865
[0342]
7TABLE 7 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N Cu Ni 3A
0.10 1.20 1.42 0.01 0.003 0.032 0.002 1.51 -- 3B 0.11 1.10 1.51
0.01 0.002 0.033 0.002 1.45 0.63 3C 0.11 1.32 1.33 0.01 0.004 0.025
0.002 1.20 0.52 3D 0.10 1.06 1.48 0.01 0.003 0.022 0.002 1.39 0.43
3E 0.09 1.25 1.36 0.01 0.004 0.029 0.002 0.22 -- 3F 0.10 1.08 1.45
0.01 0.001 0.030 0.002 0.75 -- 3G 0.11 1.15 1.52 0.01 0.002 0.033
0.002 0.96 -- 3H 0.10 1.10 1.55 0.01 0.002 0.025 0.002 1.22 0.66 3I
0.11 1.09 1.48 0.01 0.001 0.033 0.002 1.36 -- 3J 0.11 1.12 1.62
0.01 0.002 0.029 0.001 1.42 -- 3K 0.10 1.25 1.39 0.01 0.002 0.032
0.002 1.38 -- 3L 0.09 1.10 1.45 0.01 0.003 0.025 0.002 1.29 -- 3M
0.10 1.35 1.50 0.01 0.002 0.030 0.002 1.44 -- 3N 0.11 1.26 1.46
0.01 0.001 0.028 0.001 1.33 0.52 TRANSFORMATION STEEL COMPOSITION
POINT (.degree. C.) NO. Cr Mo Nb Ti V Ac1 Ac3 3A -- -- -- -- -- 725
875 3B -- 0.11 -- -- -- 715 875 3C 0.12 -- -- -- -- 725 880 3D --
-- 0.01 0.01 0.01 720 870 3E -- -- -- -- -- 730 860 3F -- -- -- --
-- 720 880 3G -- -- -- -- -- 725 875 3H -- -- -- -- -- 730 875 3I
-- 0.10 -- -- -- 725 860 3J 0.10 -- -- -- -- 730 880 3K -- -- 0.01
-- -- 720 870 3L -- -- -- 0.01 -- 725 865 3M -- -- -- -- 0.01 730
875 3N 0.12 0.11 0.01 0.01 0.01 725 865
[0343] A test piece was sampled from the resultant steel strip, and
the microstructure, tensile properties, the strain age
hardenability, and the hole expanding property were investigated,
as in Example 1. The press formability was evaluated in terms of
the elongation El (ductility), strength-elongation balance
TS.times.El, and the hole expanding ratio, as in Example 1.
[0344] (1) Microstructure
[0345] A test piece was sampled from each of the resultant steel
sheets, and the microstructure of the cross-section (section L) in
the rolling direction of the steel sheet was observed with an
optical microscope and a scanning electron microscope. The volume
ratios of the ferrite, bainite, and martensite phases in the steel
sheet were determined, as in Example 1, by image analysis using a
photograph of the cross-sectional structure at a magnification of
1,000. The amount of the retained austenite was determined, as in
Example 1, by polishing the steel sheet to the central plane in the
thickness direction and by measuring diffraction X-ray intensities
at the central plane. The incident X-ray, the planes of the ferrite
phase, and the planes of retained austenite used were the same as
those in Example 1.
[0346] (2) Tensile Properties
[0347] JIS No. 5 tensile test pieces were sampled from the
resultant steel strips in the direction perpendicular to the
rolling direction, and a tensile test was carried out, as in
Example 1, in accordance with JIS Z 2241 to determine yield
strength YS, tensile strength TS, and elongation El.
[0348] (3) Strain Age Hardenability
[0349] JIS No. 5 test pieces were sampled in the direction
perpendicular to the rolling direction from the resultant steel
strips (cold-rolled annealed sheets). A plastic deformation of 5%
was applied as a pre-deformation (tensile prestrain), as in Example
1. After a heat treatment at 250.degree. C. for 20 minutes, a
tensile test was carried out to determine tensile properties (yield
stress YS.sub.HT, and tensile strength TS.sub.HT) and to calculate
.DELTA.YS=YS.sub.HT-YS, and .DELTA.TS=TS.sub.HT-TS, wherein
YS.sub.HT and TS.sub.HT were yield stress and tensile strength
after the pre-deformation -heat treatment, and YS and TS were yield
stress and tensile strength of the steel strips (cold-rolled
annealed sheets).
[0350] (4) Hole Expanding Property
[0351] A hole was formed by punching a test piece sampled from the
resultant steel strip in accordance with Japan Iron and Steel
Federation Standard JFS T 1001-1996 with a punch having a diameter
of 10 mm. Then, the hole was expanded with a conical punch having a
vertical angle of 60.degree. so that burrs were produced on the
outside until cracks passing through the thickness form, thereby
determining the hole expanding ratio .lambda., as in Example 1.
[0352] The results are shown in Table 9.
8 TABLE 9 MICROSTRUCTURE SECONDARY PHASE RETAINED COLD-ROLLED SHEET
PROPERTIES STEEL FERRITE AUSTENITE VOLUME TENSILE PROPERTIES SHEET
STEEL VOLUME RATIO VOLUME RATIO YS TS EI NO. NO. (%) KIND RATIO %
(%) (MPa) (MPa) (%) TS .times. EI 3-1 3A 90 A, B 6 10 475 630 34
21420 3-2 3B 92 A, B 5 8 500 660 32 21120 3-3 0 P, B, M 0 100 690
730 11 8030 3-4 100 -- 0 0 650 670 11 7370 3-5 3C 92 A, B 5 8 490
650 33 21450 3-6 3D 91 A, B 5 9 500 670 32 21440 3-7 3E 93 A, B 3 7
400 530 40 21200 3-8 3F 94 A, B 4 6 450 590 36 21240 3-9 3G 93 A, B
5 7 460 610 35 21350 3-10 3H 90 A, B 6 10 465 620 34 21080 3-11 3I
92 A, B 5 8 460 610 34 20740 3-12 3J 90 A, B 6 10 500 660 32 21120
3-13 3K 92 A, B 6 8 480 640 33 21120 3-14 3L 91 A, B 5 9 470 630 33
20790 3-15 3M 90 A, B 5 10 475 630 34 21420 3-15 3N 92 A, B 4 8 460
610 34 20740 3-17 3A 90 P 0 10 510 600 28 16800 3-18 91 B 0 9 540
630 25 15750 3-19 90 M 0 10 420 650 27 17550 3-20 92 M 0 8 430 640
28 17920 PROPERTIES AFTER STRAIN AGE HOLE PRE-DEFORMATION -
HARDENING EXPANSION STEEL HEAT TREATMENT PROPERTIES HOLE SHEET
YS.sub.HT TS.sub.HT .DELTA.YS .DELTA.TS EXPANDING NO. (MPa) (MPa)
(MPa) (MPa) RATIO .lambda. % REMARKS 3-1 710 790 235 160 140
EXAMPLE 3-2 750 830 250 170 135 EXAMPLE 3-3 740 760 50 30 60 COMP.
EX. 3-4 690 695 40 25 130 COMP. EX. 3-5 730 810 240 160 135 EXAMPLE
3-6 750 825 250 155 130 EXAMPLE 3-7 500 550 100 20 50 COMP. EX. 3-8
670 740 220 150 145 EXAMPLE 3-9 690 765 230 155 140 EXAMPLE 3-10
700 780 235 160 130 EXAMPLE 3-11 705 780 245 170 135 EXAMPLE 3-12
740 820 240 160 130 EXAMPLE 3-13 730 810 250 170 130 EXAMPLE 3-14
720 795 250 165 135 EXAMPLE 3-15 715 790 240 160 140 EXAMPLE 3-15
705 780 245 170 130 EXAMPLE 3-17 590 650 80 50 70 COMP. EX. 3-18
605 670 65 40 120 COMP. EX. 3-19 725 805 305 155 125 COMP. EX. 3-20
720 800 290 160 120 COMP. EX. F: FERRITE, A: RETAINED AUSTENITE, M:
MARTENSITE, P: PEARLITE, B: BAINITE
[0353] All Examples according to the present invention are
cold-rolled steel sheets having a high elongation El, a high
strength-elongation balance TS.times.El, a high hole expanding
ratio .lambda., and excellent press formability including stretch
flanging formability. In addition, Examples according to the
present invention each show a very large .DELTA.TS, suggesting that
the samples have excellent strain age hardenability. Comparative
Examples outside the scope of the present invention, in contrast,
suggest that the samples each have a low elongation El, a low
TS.times.El, a small hole expanding ratio .lambda., a low
.DELTA.TS, and decreased press formability and strain age
hardenability.
Example 4
[0354] Molten steels having the compositions shown in Table 10 were
made in a converter and cast into steel slabs by a continuous
casting process. Each of these steel slabs were reheated to
1,250.degree. C., and hot-rolled by a hot rolling step of hot
rolling with a finish rolling end temperature of 900.degree. C. and
a coiling temperature of 600.degree. C. into a hot-rolled steel
strip (hot-rolled sheet) having a thickness of 4.0 mm. Then, the
hot-rolled steel strip (hot-rolled sheet) was subjected to a cold
rolling step of pickling and cold-rolling into a cold rolled steel
strip (cold-rolled sheet) having a thickness of 1.2 mm. Thereafter,
the cold-rolled steel strip (cold-rolled sheet) was subjected to
recrystallization annealing step comprising a heating and soaking
treatment and a subsequent retaining treatment under the conditions
shown in Table 11 on a continuous annealing line to obtain
cold-rolled annealed sheet. The resultant steel strip (cold-rolled
annealed sheet) was further temper-rolled at an reduction of
0.8%.
9TABLE 10 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N Mo Cr 4A
0.10 1.21 1.45 0.01 0.003 0.032 0.002 0.45 0.15 4B 0.11 1.12 1.52
0.01 0.002 0.032 0.002 0.32 -- 4C 0.11 1.30 1.35 0.01 0.003 0.028
0.002 0.48 -- 4D 0.10 1.05 1.50 0.01 0.004 0.033 0.002 -- -- 4E
0.09 1.26 1.38 0.01 0.004 0.032 0.002 0.35 -- 4F 0.10 1.10 1.48
0.01 0.003 0.031 0.002 -- 0.50 4G 0.11 1.16 1.53 0.01 0.004 0.032
0.002 -- -- 4H 0.12 1.20 1.52 0.01 0.002 0.028 0.002 0.35 -- 4I
0.10 1.18 1.45 0.01 0.002 0.030 0.002 -- 0.25 4J 0.11 1.10 1.36
0.01 0.003 0.031 0.002 0.45 -- 4K 0.12 1.15 1.45 0.01 0.001 0.025
0.002 0.30 -- 4L 0.11 1.08 1.50 0.01 0.003 0.032 0.002 0.25 0.15
TRANSFORMATION STEEL COMPOSITION POINT (.degree. C.) NO. W Nb Ti V
Ac1 Ac3 4A -- -- -- -- 740 880 4B -- 0.04 -- 0.05 735 875 4C --
0.05 0.03 -- 740 885 4D 0.54 -- -- -- 735 875 4E -- -- 0.05 -- 735
880 4F -- 0.05 -- -- 730 885 4G -- -- -- -- 725 830 4H -- -- -- --
740 870 4I -- -- -- -- 735 860 4J -- -- -- -- 730 860 4K -- 0.03
0.01 0.01 735 850 4L 0.10 -- -- -- 740 865
[0355]
10 TABLE 11 RECRYSTALLIZATION ANNEALING STEP HEATING HOT ROLLING
STEP SOAKING SLAB FINISH COLD ROLLING TREATMENT RETAINING STEEL
REHEATING ROLLING END COILING. STEP HEATING TREATMENT SHEET STEEL
TEMP. TEMP. TEMP. COLD ROLLING SOAKING TEMP. RETENTION NO. NO.
(.degree. C.) FDT .degree. C. CT .degree. C. REDUCTION % TEMP.
(.degree. C.) (.degree. C.) TIME(s) 4-1 4A 1250 900 600 70 800 400
300 4-2 4B 1250 900 600 70 800 400 300 4-3 1250 900 600 70 980 --
-- 4-4 1250 900 600 70 680 400 300 4-5 4C 1250 900 600 70 800 400
300 4-6 4D 1250 900 600 70 800 400 300 4-7 4E 1250 900 600 70 800
400 300 4-8 4F 1250 900 600 70 800 400 300 4-9 4G 1250 900 600 70
800 400 300 4-10 4H 1250 900 600 70 800 400 300 4-11 4I 1250 900
600 70 800 400 300 4-12 4J 1250 900 600 70 800 400 300 4-13 4K 1250
900 600 70 800 400 300 4-14 4L 1250 900 600 70 800 400 300 4-15 4A
1250 900 600 70 800 250 300 4-16 1250 900 600 70 800 550 300
[0356] A test piece was sampled from the resultant steel strip, and
the microstructure, the tensile properties, the strain age
hardenability, and the hole expanding property were investigated,
as in Example 3.
[0357] The results are shown in Table 12.
11 TABLE 12 MICROSTRUCTURE SECONDARY PHASE FERRITE RETAINED
COLD-ROLLED SHEET PROPERTIES STEEL VOLUME AUSTENITE VOLUME TENSILE
PROPERTIES SHEET STEEL RATIO VOLUME RATIO YS TS NO. NO. (%) KIND
RATIO % (%) (MPa) (MPa) El (%) TS .times. El 4-1 4A 91 A, B 6 9 470
630 34 21420 4-2 4B 92 A, B 5 8 500 660 32 21120 4-3 0 P, B, M 0
100 560 740 12 8880 4-4 100 -- 0 0 500 660 11 7260 4-5 4C 92 A, B 5
8 480 640 33 21120 4-6 4D 94 A, B 4 6 470 630 34 21420 4-7 4E 92 A,
B 5 8 490 650 33 21450 4-8 4F 93 A, B 4 7 470 620 34 21080 4-9 4G
94 A, B 3 6 460 620 34 21080 4-10 4H 92 A, B 5 8 475 630 33 20790
4-11 4I 90 A, B 4 10 480 640 33 21120 4-12 4J 91 A, B 5 9 485 650
32 20800 4-13 4K 92 A, B 4 8 470 630 34 21420 4-14 4L 90 A, B 5 10
465 620 34 21080 4-15 4A 93 M 0 7 380 630 28 17640 4-16 92 P 0 8
550 650 24 15600 PROPERTIES AFTER STRAIN AGE HOLE PRE-DEFORMATION -
HARDENING EXPANSION STEEL HEAT TREATMENT PROPERTIES HOLE SHEET
YS.sub.HT TS.sub.HT .DELTA.YS .DELTA.TS EXPANDING NO. (MPa) (MPa)
(MPa) (MPa) RATIO .lambda. % REMARKS 4-1 700 780 230 150 140
EXAMPLE 4-2 740 820 240 160 130 EXAMPLE 4-3 680 760 120 20 60 COMP.
EX. 4-4 610 675 110 15 130 COMP. EX. 4-5 710 790 230 150 120
EXAMPLE 4-6 700 775 230 145 130 EXAMPLE 4-7 720 800 230 150 120
EXAMPLE 4-8 680 760 210 140 120 EXAMPLE 4-9 570 630 110 10 60 COMP.
EX. 4-10 710 790 235 160 130 EXAMPLE 4-11 725 805 245 165 120
EXAMPLE 4-12 730 810 245 160 120 EXAMPLE 4-13 710 790 240 160 130
EXAMPLE 4-14 700 775 235 155 120 EXAMPLE 4-15 710 790 330 160 110
COMP. EX. 4-16 620 680 70 30 70 COMP. EX. F: FERRITE, A: RETAINED
AUSTENITE, M: MARTENSITE, P: PEARLITE, B: BAINITE
[0358] All Examples according to the present invention show a high
elongation El, a high strength-ductility balance TS.times.El, and a
high hole expanding ratio .lambda., suggesting that the samples
have excellent press formability including stretch flanging
formability. In addition, Examples according to the present
invention show a very large .DELTA.TS, suggesting that the samples
have excellent strain age hardenability. Comparative Examples
outside the scope of the present invention, in contrast, suggest
that the samples have a low elongation El, a low TS.times.El, a
small hole expanding ratio .lambda., a low .DELTA.TS, and decreased
press formability and strain age hardenability.
Example 5
[0359] Molten steels having the compositions shown in Table 13 were
made in a converter and cast into steel slabs by a continuous
casting process. These slabs were hot-rolled under the conditions
shown in Table 14 into hot-rolled steel strips (hot-rolled
sheets).
[0360] After pickling, each of these hot-rolled steel strips
(hot-rolled sheets) was subjected to a primary heat treatment step
on a continuous annealing line (CAL) under the conditions shown in
Table 14 and a secondary heat treatment step on a continuous
hot-dip galvanizing line (CGL) under the conditions shown in Table.
14. Then, the sheet was subjected to a hot-dip galvanizing
treatment step of performing a hot-dip galvanizing which forms a
hot-dip galvanizing layer on the surfaces of the steel sheet. Then,
an alloying treatment step of alloying the hot-dip galvanizing
layer was applied under the conditions shown in Table 14. Some of
the steel sheets were left as hot-dip galvanized.
[0361] After further pickling, the hot-rolled steel strip
(hot-rolled sheet) obtained by the above-mentioned hot rolling was
subjected to a cold rolling step under the conditions shown in
Table 14 into a cold-rolled steel strip (cold-rolled sheet). Then,
the cold-rolled steel strip (cold-rolled sheet) was subjected to a
primary heat treatment step on a continuous annealing line (CAL)
under the conditions shown in Table 14. After a secondary heat
treatment step on the continuous hot-dip galvanizing line (CGL)
under the conditions shown in Table 14, a hot-dip galvanizing
treatment step was performed. Then, an alloying treatment step was
performed under the conditions shown in Table 14. Some of the steel
sheets were left as hot-dip galvanized.
[0362] Prior to the secondary heat treatment-step on the continuous
hot-dip galvanizing line (CGL), some of the steel sheets after the
primary heat treatment step were subjected to a pickling treatment
shown in Table 14. The pickling treatment was carried out in a
pickling bath on the entry side of the CGL.
[0363] The galvanizing bath temperature was within the range of 460
to 480.degree. C., and the temperature of the steel sheet to be
dipped was within the range of the galvanizing bath temperature to
(bath temperature +10.degree. C.). In the alloying treatment, the
sheet was reheated within the temperature range of 480 to
540.degree. C., and held at the temperature for 15 to 28 seconds.
The cooling rate after the alloying treatment was 10.degree.
C./second. The plated steel sheet was further temper rolled at a
reduction of 1.0%.
12TABLE 13 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N Cu Ni 5A
0.08 0.72 2.05 0.01 0.003 0.032 0.002 1.48 -- 5B 0.07 0.52 2.22
0.01 0.001 0.033 0.002 1.44 0.62 5C 0.09 0.77 1.85 0.01 0.004 0.028
0.002 1.28 0.55 5D 0.08 0.65 1.95 0.01 0.005 0.032 0.002 1.33 0.42
5E 0.07 0.55 2.05 0.01 0.004 0.033 0.002 0.14 -- 5F 0.08 0.70 2.22
0.01 0.003 0.033 0.002 0.72 -- 5G 0.07 0.68 1.85 0.01 0.005 0.036
0.002 0.95 -- 5H 0.08 0.77 2.05 0.01 0.003 0.032 0.002 1.45 0.75 5I
0.09 0.80 1.85 0.01 0.002 0.028 0.002 1.29 -- 5J 0.07 0.75 2.05
0.01 0.005 0.030 0.002 1.38 -- 5K 0.08 0.68 1.95 0.01 0.003 0.025
0.002 1.40 -- 5L 0.07 0.70 2.10 0.01 0.004 0.030 0.002 1.35 -- 5M
0.08 0.75 1.80 0.01 0.002 0.031 0.002 1.25 -- 5N 0.09 0.68 2.00
0.01 0.003 0.035 0.002 1.35 0.60 TRANSFORMATION STEEL COMPOSITION
(wt. %) POINT (.degree. C.) NO. Cr, Mo Nb, Ti, V Ac1 Ac3 5A -- --
715 875 5B Mo: 0.15 -- 720 870 5C Cr: 0.15 -- 725 875 5D -- Nb:
0.01, 715 870 Ti: 0.01, V: 0.01 5E -- -- 715 875 5F -- -- 715 870
5G -- -- 715 875 5H -- -- 715 870 5I Cr: 0.12 -- 720 875 5J Mo:
0.15 -- 715 870 5K -- Nb: 0.01 720 875 5L -- Ti: 0.01 715 870 5M --
V: 0.01 725 870 5N Cr: 0.13, Nb: 0.01, 710 875 Mo: 0.15 V: 0.01
[0364]
13 TABLE 14 HOT ROLLING STEP COLD ROLLING STEP PRIMARY HEAT SLAB
FINISH FINAL COLD FINAL TREATMENT STEP STEEL REHEATING ROLLING
COILING THICK- ROLLING THICK- COOLING SHEET STEEL TEMP. END TEMP.
TEMP. NESS REDUCTION NESS HEATING RATE PICKLING NO. NO. (.degree.
C.) FDT .degree. C. CT .degree. C. mm % mm LINE TEMP. .degree. C.
.degree. C./s TREATMENT 5-1 5A 1250 850 600 1.2 -- -- CAL 880 20
YES 5-2 5B 1250 850 600 1.2 -- -- CAL 880 20 -- 5-3 YES 5-4 5-5 5-6
5C 1250 850 600 1.2 -- -- CAL 880 20 YES 5-7 5D 1250 850 600 1.2 --
-- CAL 880 20 YES 5-8 5E 1250 850 600 1.2 -- -- CAL 880 20 YES 5-9
5F 1250 850 600 1.2 -- -- CAL 880 20 YES 5-10 5G 1250 850 600 1.2
-- -- CAL 880 20 YES 5-11 5A 1250 850 600 4.0 70 1.2 CAL 880 20 YES
5-12 5B 1250 850 600 4.0 70 1.2 CAL 880 20 -- 5-13 CAL 880 20 YES
5-14 CAL 880 20 YES 5-15 CAL 880 20 YES 5-16 5C 1250 850 600 4.0 70
1.2 CAL 880 20 YES 5-17 5D 1250 850 600 4.0 70 1.2 CAL 880 20 YES
5-18 5E 1250 850 600 4.0 70 1.2 CAL 880 20 YES 5-19 5F 1250 850 600
4.0 70 1.2 CAL 880 20 YES 5-20 5G 1250 850 600 4.0 70 1.2 CAL 880
20 YES 5-21 5H 1250 850 600 4.0 70 1.2 CAL 880 20 YES 5-22 5I 1250
850 600 4.0 70 1.2 CAL 880 20 YES 5-23 5J 1250 850 600 4.0 70 1.2
CAL 880 20 YES 5-24 5K 1250 850 600 4.0 70 1.2 CAL 880 20 YES 5-25
5L 1250 850 600 4.0 70 1.2 CAL 880 20 YES 5-26 5M 1250 850 600 4.0
70 1.2 CAL 880 20 YES 5-27 5N 1250 850 600 4.0 70 1.2 CAL 880 20
YES HOT-DIP SECONDARY HEAT GALVANIZING TREATMENT STEP COOLING
TEMPER STEEL KIND KIND RATE AFTER ALLOYING TREATMENT ROLLING SHEET
OF HEATING COOLING OF GALVANIZING STEP REDUCTION NO. LINE TEMP.
.degree. C. RATE* .degree. C./s LINE **.degree. C./s TEMP. .degree.
C. % 5-1 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-2 CGL 800 20 CGL 10
ALLOYING 500 1.0 5-3 CGL 780 20 CGL 10 ALLOYING 500 1.0 5-4 CGL 980
20 CGL 10 ALLOYING 500 1.0 5-5 CGL 650 20 CGL 10 ALLOYING 500 1.0
5-6 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-7 CGL 820 20 CGL 10
ALLOYING 500 1.0 5-8 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-9 CGL 780
20 CGL 10 NON- -- 1.0 ALLOYING 5-10 CGL 800 20 CGL 10 ALLOYING 500
1.0 5-11 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-12 CGL 820 20 CGL 10
ALLOYING 500 1.0 5-13 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-14 CGL
980 20 CGL 10 ALLOYING 500 1.0 5-15 CGL 680 20 CGL 10 ALLOYING 500
1.0 5-16 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-17 CGL 800 20 CGL 10
NON- -- 1.0 ALLOYING 5-18 CGL 780 20 CGL 10 ALLOYING 500 1.0 5-19
CGL 800 20 CGL 10 ALLOYING 500 1.0 5-20 CGL 820 20 CGL 10 ALLOYING
500 1.0 5-21 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-22 CGL 800 20 CGL
10 ALLOYING 500 1.0 5-23 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-24
CGL 800 20 CGL 10 ALLOYING 500 1.0 5-25 CGL 800 20 CGL 10 ALLOYING
500 1.0 5-26 CGL 800 20 CGL 10 ALLOYING 500 1.0 5-27 CGL 800 20 CGL
10 ALLOYING 500 1.0 *COOLING RATE UNTIL 480.degree. C. **COOLING
RATE UNTIL 300.degree. C.
[0365] For the hot-dip galvanized steel sheet (steel strip)
obtained through the above-mentioned steps, the microstructure, the
tensile properties, the strain age hardenability, and the hole
expanding ratio were determined, as in Example 1. Press formability
was evaluated in terms of elongation El (ductility), and hole
expanding ratio.
[0366] (1) Microstructure
[0367] The microstructure of the cross-section (section L) in the
rolling direction of the steel sheet was observed with an optical
microscope and a scanning electron microscope. The volume ratios of
the ferrite phase, lath martensite phase, tempered martensite
phase, and martensite phase were determined, as in Example 1, by
image analysis using a photograph of cross-sectional structure at a
magnification of 1,000. The amount of retained austenite was
determined, as in Example 1, by polishing the steel sheet to the
central plane in the thickness direction and by measuring
diffraction X-ray intensities at the central plane. The incident
X-ray, the planes of the ferrite phase, and the planes of retained
austenite used were the same as those in Example 1.
[0368] (2) Tensile properties
[0369] JIS No. 5 tensile test pieces were sampled from the
resultant steel strips in the direction perpendicular to the
rolling direction, and a tensile test was carried out in accordance
with JIS Z 2241 to determine the yield strength YS, the tensile
strength TS, and the elongation El, as, in Example 1.
[0370] (3) Strain Age Hardenability
[0371] JIS No. 5 test pieces were sampled from the resultant steel
strips in the direction perpendicular to the rolling direction, and
a plastic deformation of 5% was applied as a pre-deformation
(tensile prestrain), as in Example 1. After a heat treatment at
250.degree. C. for 20 minutes, a tensile test was carried out to
determine tensile properties (yield stress YS.sub.TH, and tensile
strength TS.sub.HT) and to calculate .DELTA.YS=YS.sub.TH-YS, and
.DELTA.TS=TS.sub.HT-TS, wherein YS.sub.TS and TS.sub.HT were yield
stress and tensile strength after the pre-deformation--heat
treatment, and YS and TS were yield stress and tensile strength of
the steel strips.
[0372] (4) Hole Expanding Ratio
[0373] A hole was formed by punching a test piece sampled from the
resultant steel strip in accordance with Japan Iron and Steel
Federation Standard JFS T 1001-1996 with a punch having a diameter
of 10 mm. Then, the hole was expanded with a conical punch having a
vertical angle of 60.degree. C. so that burrs were produced on the
outside until cracks passing through the thickness form, thereby
determining the hole expanding ratio .lambda., as in Example 1.
[0374] The results are shown in Table 15.
14 TABLE 15 MICROSTRUCTURE PRIMARY PHASE SECONDARY PHASE TEMPERED
RETAINED PLATED SHEET PROPERTIES STEEL FERRITE MARTENSITE AUSTENITE
TENSILE PROPERTIES SHEET STEEL VOLUME VOLUME VOLUME VOLUME VOLUME
YS TS TS .times. El NO. NO. RATIO % RATIO % RATIO % KIND* RATIO %
RATIO % (MPa) (MPa) El (%) (MPa %) 5-1 5A 57 35 92 A, B, 5 8 470
620 34 21080 5-2 5B 52 40 92 A, B, 4 8 480 640 33 21120 5-3 51 40
91 A, B, 5 9 470 620 34 21080 5-4 0 0 0 M, P, B 0 100 670 710 11
7810 5-5 60 40 100 -- 0 0 620 650 12 7800 5-6 5C 58 35 93 A, B 4 7
470 630 34 21420 5-7 5D 57 35 92 A, B 5 8 490 650 33 21450 5-8 5E
53 40 93 A, B 7 7 380 510 42 21420 5-9 5F 37 55 92 A, B 4 8 430 570
37 21090 5-10 5G 53 40 93 A, B 5 7 450 590 36 21240 5-11 5A 57 35
92 A, B 7 8 470 630 34 21420 5-12 5B 52 40 92 A, B 5 8 500 660 32
21120 5-13 53 40 93 A, B 6 7 480 640 33 21120 5-14 0 0 0 M, P, B 0
100 680 720 12 8640 5-15 65 35 100 -- 0 0 620 660 11 7260 5-16 5C
52 40 92 A, B 4 8 490 650 33 21450 5-17 5D 53 40 93 A, B 5 7 500
660 32 21120 5-18 5E 48 45 93 A, B 4 7 390 520 41 21320 5-19 5F 44
50 94 A, B 5 6 440 580 37 21460 5-20 5G 57 35 92 A, B 5 8 450 600
35 21000 5-21 5H 51 40 91 A, B 5 9 445 590 35 20650 5-22 5I 55 35
90 A, B 5 10 460 610 34 20740 5-23 5J 52 40 92 A, B 4 8 450 600 35
21000 5-24 5K 53 40 93 A, B 5 7 470 620 34 21080 5-25 5L 56 35 91
A, B 6 9 475 630 33 20790 5-26 5M 60 30 90 A, B 5 10 460 610 34
20740 5-27 5N 52 40 92 A, B 4 8 455 600 35 21000 HOLE PROPERTIES
AFTER STRAIN AGE EXPANSION STEEL PREDEFORMATION - HARDENING HOLE
SHEET HEAT TREATMENT PROPERTIES EXPANDING NO. YS.sub.HT (MPa)
TS.sub.HT (MPa) .DELTA.YS (MPa) .DELTA.TS (MPa) RATIO .lambda. %
REMARKS 5-1 700 775 230 155 140 EXAMPLE 5-2 725 805 245 165 135
EXAMPLE 5-3 710 785 240 165 135 EXAMPLE 5-4 710 740 40 30 65 COMP.
EX. 5-5 650 675 30 25 130 COMP. EX. 5-6 710 785 240 155 135 EXAMPLE
5-7 725 805 235 155 130 EXAMPLE 5-8 480 530 100 20 60 COMP. EX. 5-9
650 720 220 150 140 EXAMPLE 5-10 675 745 225 155 135 EXAMPLE 5-11
715 790 245 160 145 EXAMPLE 5-12 750 830 250 170 140 EXAMPLE 5-13
730 810 250 170 140 EXAMPLE 5-14 720 750 40 30 70 COMP. EX. 5-15
650 685 30 25 60 COMP. EX. 5-16 730 810 240 160 140 EXAMPLE 5-17
735 815 235 155 135 EXAMPLE 5-18 490 540 100 20 60 COMP. EX. 5-19
655 725 215 145 135 EXAMPLE 5-20 675 750 225 150 140 EXAMPLE 5-21
680 755 235 165 130 EXAMPLE 5-22 695 770 235 160 135 EXAMPLE 5-23
680 755 230 155 130 EXAMPLE 5-24 710 780 240 160 130 EXAMPLE 5-25
720 795 245 165 135 EXAMPLE 5-26 695 770 235 160 130 EXAMPLE 5-27
680 755 225 155 130 EXAMPLE *M: MARTENSITE, P: PEARLITE, B:
BAINITE, A: RETAINED AUSTENITE
[0375] All Examples according to the present invention each show a
high elongation El and a high hole expanding ratio .lambda.,
suggesting that the samples are hot-dip galvanized steel sheets
having an excellent stretch flanging formability. In addition,
Examples according to the present invention showed a very large
.DELTA.TS, suggesting that the samples are steel sheets having
excellent strain age hardenability. Comparative Examples outside
the scope of the invention, in contrast, suggest that the samples
are steel sheets having a low elongation El, a small hole expanding
ratio .lambda., a low .DELTA.TS, and decreased press formability
and strain age hardenability.
Example 6
[0376] Molten steels having the compositions shown in Table 16 was
made in a converter and cast into steel slabs by a continuous
casting process. Each of these steel slabs were reheated to
1,250.degree. C., and hot-rolled by a hot rolling step of hot
rolling with a finish rolling end temperature of 900.degree. C. and
a coiling temperature of 600.degree. C. into hot-rolled steel strip
(hot-rolled sheet) having a thickness of 4.0 mm. Then, the
hot-rolled steel strip (hot-rolled sheet) was subjected to a cold
rolling step of pickling and cold-rolling into cold-rolled steel
strip (cold-rolled sheet) having a thickness of 1.2 mm. Then, the
cold-rolled steel strip (cold-rolled sheet) was subjected to a
primary heat treatment step on a continuous annealing line (CAL)
under the conditions shown in Table 17. Then, the sheet was
subjected to a secondary heat treatment step on a continuous
hot-dip galvanizing line (CGL) under the conditions shown in Table
17 and then, subjected to a hot-dip galvanizing treatment step to
form a hot-dip galvanizing layer on the surfaces of the steel
sheet. In addition, an alloying treatment step was applied under
the conditions shown in FIG. 17. The cooling rate after the
alloying treatment was 10.degree. C./second. Some of the steel
strips (steel sheets) were left as hot-dip galvanized.
15TABLE 16 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N 6A 0.07
0.77 2.00 0.01 0.003 0.033 0.002 6B 0.08 0.55 2.22 0.01 0.001 0.033
0.002 6C 0.08 0.75 1.80 0.01 0.004 0.020 0.002 6D 0.09 0.63 1.98
0.01 0.005 0.025 0.002 6E 0.07 0.65 2.02 0.01 0.003 0.033 0.002 6F
0.08 0.70 1.90 0.01 0.005 0.035 0.002 6G 0.07 0.58 2.08 0.01 0.004
0.032 0.002 6H 0.08 0.75 2.22 0.01 0.004 0.022 0.002 6I 0.08 0.77
1.98 0.01 0.003 0.032 0.002 6J 0.07 0.68 2.05 0.01 0.002 0.035
0.002 6K 0.09 0.70 1.98 0.01 0.001 0.028 0.002 TRANSFORMATION STEEL
COMPOSITION (wt. %) POINT (.degree. C.) NO. Cr, Mo, W Nb, Ti, V Ac1
Ac3 6A Cr: 0.20, -- 715 870 Mo: 0.43 6B Mo: 0.33 Nb: 0.04, 720 865
V: 0.05 6C Mo: 0.48 Nb: 0.05, 725 880 Ti: 0.03 6D W: 0.54 -- 715
865 6E Mo: 0.36 Ti: 0.05 715 875 6F Cr: 0.50 Nb: 0.05 715 865 6G --
-- 715 865 6H Mo: 0.35 -- 715 870 6I Cr: 0.25 -- 710 860 6J Mo:
0.15, -- 720 865 Cr: 0.10, W: 0.11 6K Mo: 0.25, V: 0.05 715 865 Cr:
0.10
[0377]
16 TABLE 17 HOT ROLLING COLD STEP ROLLING FINISH STEP PRIMARY HEAT
SLAB ROLLING FINAL COLD FINAL TREATMENT STEP STEEL REHEATING END
COILING THICK- ROLLING THICK- HEATING COOLING SHEET STEEL TEMP.
TEMP. TEMP. NESS REDUCTION NESS TEMP. RATE PICKLING NO. NO.
(.degree. C.) FDT .degree. C. CT .degree. C. mm % mm LINE .degree.
C. .degree. C./s TREATMENT 6-1 6A 1250 850 600 1.2 -- -- CAL 880 20
YES 6-2 6B 1250 850 600 1.2 -- -- CAL 880 20 -- 6-3 YES 6-4 6-5 6-6
6C 1250 850 600 1.2 -- -- CAL 880 20 YES 6-7 6D 1250 850 600 1.2 --
-- CAL 880 20 YES 6-8 6E 1250 850 600 1.2 -- -- CAL 880 20 YES 6-9
6F 1250 850 600 1.2 -- -- CAL 880 20 YES 6-10 6G 1250 850 600 1.2
-- -- CAL 880 20 YES 6-11 6A 1250 850 600 4.0 70 1.2 CAL 880 20 YES
6-12 6B 1250 850 600 4.0 70 1.2 CAL 880 20 -- 6-13 CAL 880 20 YES
6-14 CAL 880 20 YES 6-15 CAL 880 20 YES 6-16 6C 1250 850 600 4.0 70
1.2 CAL 880 20 YES 6-17 6D 1250 850 600 4.0 70 1.2 CAL 880 20 YES
6-18 6E 1250 850 600 4.0 70 1.2 CAL 880 20 YES 6-19 6F 1250 850 600
4.0 70 1.2 CAL 880 20 YES 6-20 6G 1250 850 600 4.0 70 1.2 CAL 880
20 YES 6-21 6H 1250 850 600 4.0 70 1.2 CAL 880 20 YES 6-22 6I 1250
850 600 4.0 70 1.2 CAL 880 20 YES 6-23 6J 1250 850 600 4.0 70 1.2
CAL 880 20 YES 6-24 6K 1250 850 600 4.0 70 1.2 CAL 880 20 YES
HOT-DIP SECONDARY HEAT GALVANIZING TREATMENT STEP COOLING ALLOYING
TEMPER STEEL KIND KIND RATE AFTER TREATMENT ROLLING SHEET OF
HEATING COOLING OF GALVANIZING STEP REDUCTION NO. LINE TEMP.
.degree. C. RATE* .degree. C./s LINE **.degree. C./s TEMP. .degree.
C. % 6-1 CGL 780 20 CGL 10 ALLOYING 500 1.0 6-2 CGL 800 20 CGL 10
ALLOYING 500 1.0 6-3 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-4 CGL 980
20 CGL 10 ALLOYING 500 1.0 6-5 CGL 650 20 CGL 10 ALLOYING 500 1.0
6-6 CGL 780 20 CGL 10 ALLOYING 500 1.0 6-7 CGL 820 20 CGL 10
ALLOYING 500 1.0 6-8 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-9 CGL 800
20 CGL 10 NON- -- 1.0 ALLOYING 6-10 CGL 800 20 CGL 10 ALLOYING 500
1.0 6-11 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-12 CGL 820 20 CGL 10
ALLOYING 500 1.0 6-13 CGL 780 20 CGL 10 ALLOYING 500 1.0 6-14 CGL
980 20 CGL 10 ALLOYING 500 1.0 6-15 CGL 680 20 CGL 10 ALLOYING 500
1.0 6-16 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-17 CGL 800 20 CGL 10
NON- -- 1.0 ALLOYING 6-18 CGL 780 20 CGL 10 ALLOYING 500 1.0 6-19
CGL 800 20 CGL 10 ALLOYING 500 1.0 6-20 CGL 820 20 CGL 10 ALLOYING
500 1.0 6-21 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-22 CGL 800 20 CGL
10 ALLOYING 500 1.0 6-23 CGL 800 20 CGL 10 ALLOYING 500 1.0 6-24
CGL 800 20 CGL 10 ALLOYING 500 1.0 *COOLING RATE UNTIL 480.degree.
C. **COOLING RATE UNTIL 300.degree. C.
[0378] A piece was sampled from the resultant hot-dip galvanized
steel strip, and the microstructure, the tensile properties, the
strain age hardenability, and the bore expanding property were
investigated, as in Example 5.
[0379] The results are shown in Table 18.
17 TABLE 18 MICROSTRUCTURE PRIMARY PHASE SECONDARY PHASE TEMPERED
RETAINED PLATED SHEET PROPERTIES STEEL FERRITE MARTENSITE AUSTENITE
TENSILE PROPERTIES SHEET STEEL VOLUME VOLUME VOLUME VOLUME VOLUME
YS TS El TS .times. El NO. NO. RATIO % RATIO % RATIO % KIND* RATIO
% RATIO % (MPa) (MPa) (%) (MPa %) 6-1 6A 56 35 91 A, B 6 9 460 610
35 21350 6-2 6B 52 40 92 A, B 5 8 475 630 34 21420 6-3 50 40 90 A,
B 6 10 460 610 35 21350 6-4 0 0 O M, P, B 0 100 660 700 11 7700 6-5
60 40 100 -- 0 0 620 660 12 7920 6-6 6C 47 45 92 A, B 5 8 570 620
34 21080 6-7 6D 53 40 93 A, B 5 7 480 640 33 21120 6-8 6E 57 35 92
A, B 6 8 390 520 41 21320 6-9 6F 48 45 93 A, B 5 7 420 560 38 21280
6-10 6G 53 40 93 A, B 5 7 450 590 36 21240 6-11 6A 53 40 93 A, B 5
7 465 620 34 21080 6-12 6B 52 40 92 A, B 5 8 490 650 33 21450 6-13
57 35 92 A, B 5 8 475 630 34 21420 6-14 0 0 O M, P, B 0 100 650 710
12 8520 6-15 60 40 100 -- 0 0 610 650 11 7150 6-16 6C 53 40 93 A, B
5 7 480 640 33 21120 6-17 6D 62 30 92 A, B 5 8 490 650 33 21450
6-18 6E 53 40 93 A, B 4 7 390 520 41 21320 6-19 6F 49 45 94 A, B 4
6 450 590 36 21240 6-20 6G 42 50 92 A, B 5 8 460 610 35 21350 6-21
6H 36 55 91 A, B 5 9 470 630 34 21420 6-22 6I 40 50 90 A, B 4 10
465 620 34 21080 6-23 6J 50 40 90 A, B 5 10 480 640 33 21120 6-24
6K 51 40 91 A, B 5 9 470 620 34 21080 HOLE PROPERTIES AFTER STRAIN
AGE EXPANSION STEEL PREDEFORMATION - HARDENING HOLE SHEET HEAT
TREATMENT PROPERTIES EXPANDING NO. YS.sub.HT MPa TS.sub.HT MPa
.DELTA.YS MPa .DELTA.TS MPa RATIO .lambda. % REMARKS 6-1 705 780
245 170 140 EXAMPLE 6-2 730 810 255 180 135 EXAMPLE 6-3 715 790 255
180 135 EXAMPLE 6-4 720 730 60 30 55 COMP. EX. 6-5 660 685 40 25
125 COMP. EX. 6-6 715 790 145 170 135 EXAMPLE 6-7 730 810 250 170
130 EXAMPLE 6-8 620 685 230 165 130 EXAMPLE 6-9 655 725 235 165 140
EXAMPLE 6-10 560 620 110 30 50 COMP. EX. 6-11 720 795 255 175 145
EXAMPLE 6-12 755 835 265 185 140 EXAMPLE 6-13 730 810 255 180 140
EXAMPLE 6-14 720 740 70 30 60 COMP. EX. 6-15 650 675 40 25 50 COMP.
EX. 6-16 730 810 250 170 140 EXAMPLE 6-17 740 820 250 170 135
EXAMPLE 6-18 615 680 225 160 140 EXAMPLE 6-19 675 750 225 160 135
EXAMPLE 6-20 700 775 240 165 30 COMP. EX. 6-21 710 790 240 160 120
EXAMPLE 6-22 705 785 240 165 120 EXAMPLE 6-23 720 800 240 160 130
EXAMPLE 6-24 700 775 230 155 120 EXAMPLE *M: MARTENSITE, P:
PEARLITE, B: BAINITE, A: RETAINED AUSTENITE
[0380] All Examples according to the present invention show a high
elongation El and a high bore expanding ratio .lambda., suggesting
that the examples are hot-dip galvanized steel sheets having
excellent press formability. In addition, all Examples according to
the present invention show a very large .DELTA.TS, suggesting that
the samples are steel sheets having excellent strain age
hardenability. Comparative Examples outside the scope of the
invention, in contrast, suggest that the samples are steel sheets
having a low elongation El, a low .lambda., a low .DELTA.TS, and
decreased press formability and strain age hardenability.
[0381] According to the present invention, it is possible to stably
manufacture steel sheets (hot-rolled steel sheets, cold-rolled
steel sheets and hot-dip galvanized steel sheets) in which the
tensile strength is remarkably increased through a heat treatment
applied after press forming while maintaining excellent press
formability, giving industrially remarkable effects. When applying
a steel sheet of the present invention to automotive parts, there
are available advantages of easy press forming, high and stable
parts properties after completion, and sufficient contribution to
the weight reduction of the automobile body.
* * * * *