U.S. patent application number 10/181810 was filed with the patent office on 2003-07-10 for composite structure type high tensile strength steel plate, plated plate of composite structure type high tensile strength steel and method for their production.
Invention is credited to Hanazawa, Kazuhiro, Matsuoka, Saiji, Sakata, Kei, Shimizu, Tetsuo.
Application Number | 20030129444 10/181810 |
Document ID | / |
Family ID | 27481823 |
Filed Date | 2003-07-10 |
United States Patent
Application |
20030129444 |
Kind Code |
A1 |
Matsuoka, Saiji ; et
al. |
July 10, 2003 |
Composite structure type high tensile strength steel plate, plated
plate of composite structure type high tensile strength steel and
method for their production
Abstract
The invention proposes a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability, wherein the steel
sheet has a composition comprising C: 0.01-0.08 mass %, Si: not
more than 2.0 mass %, Mn: not more than 3.0 mass %, P: not more
than 0.10 mass %, S: not more than 0.02 mass %, A1: 0.005-0.20 mass
%, N: not more than 0.02 mass % and V: 0.01-0.5 mass %, provided
that V and C satisfy a relationship of
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12, and the remainder
being Fe and inevitable impurities, and has a microstructure
consisting of a ferrite phase as a primary phase and a secondary
phase including martensite phase at an area ratio of not less than
1% to a whole of the microstructure and a high-strength dual-phase
galvanized steel sheet comprising a galvanized coating on the above
steel sheet as well as a method of producing the same.
Inventors: |
Matsuoka, Saiji; (Chiba
Pref, JP) ; Hanazawa, Kazuhiro; (Chiba Pref, JP)
; Shimizu, Tetsuo; (Okayama Pref, JP) ; Sakata,
Kei; (Chiba Pref, JP) |
Correspondence
Address: |
YOUNG & THOMPSON
745 SOUTH 23RD STREET 2ND FLOOR
ARLINGTON
VA
22202
|
Family ID: |
27481823 |
Appl. No.: |
10/181810 |
Filed: |
July 23, 2002 |
PCT Filed: |
November 27, 2001 |
PCT NO: |
PCT/JP01/10340 |
Current U.S.
Class: |
428/659 ;
148/651 |
Current CPC
Class: |
Y10T 428/12799 20150115;
C22C 38/04 20130101; C22C 38/12 20130101; C21D 2211/005 20130101;
C23C 2/40 20130101; C21D 2211/008 20130101; C23C 2/02 20130101;
C22C 38/06 20130101; C22C 38/001 20130101 |
Class at
Publication: |
428/659 ;
148/651 |
International
Class: |
C21D 008/00 |
Foreign Application Data
Date |
Code |
Application Number |
Nov 28, 2000 |
JP |
2000-361.273 |
Nov 28, 2000 |
JP |
2000-361.274 |
Nov 28, 2001 |
JP |
2001-312.687 |
Oct 10, 2001 |
JP |
2001-312.688 |
Claims
1. A high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability, characterized in that the steel sheet
has a composition comprising C: 0.01-0.08 mass %, Si: not more than
2.0 mass %, Mn: not more than 3.0 mass %, P: not more than 0.10
mass %, S: not more than 0.02 mass %, Al: 0.005-0.20 mass %, N: not
more than 0.02 mass % and V: 0.01-0.5 mass %, provided that V and C
satisfy a relationship of
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12, and the remainder
being Fe and inevitable impurities, and has a microstructure
consisting of a ferrite phase as a primary phase and a secondary
phase including martensite phase at an area ratio of not less than
1% to a whole of the microstructure.
2. A high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability, characterized in that the steel sheet
has a composition comprising C: 0.01-0.08 mass %, Si: not more than
2.0 mass %, Mn: not more than 3.0 mass %, P: not more than 0.10
mass %, S: not more than 0.02 mass %, Al: 0.005-0.20 mass %, N: not
more than 0.02 mass % and V: 0.01-0.5 mass % and further comprising
not more than 0.3 mass % in total of one or two of Nb: more than 0
mass % but not more than 0.3 mass % and Ti: more than 0 mass % but
not more than 0.3 mass %, provided that V, Nb, Ti and C satisfy a
relationship of 0.5.times.C/12.ltoreq.(V/51+2.times.Nb-
/93+2.times.Ti/48).ltoreq.3.times.C/12, and the remainder being Fe
and inevitable impurities, and has a microstructure consisting of a
ferrite phase as a primary phase and a secondary phase including
martensite phase at an area ratio of not less than 1% to a whole of
the microstructure.
3. A high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to claim 2, wherein the steel
sheet comprises not more than 0.3 mass % in total of one or two of
Nb: 0.001-0.3 mass % and Ti: 0.001-0.3 mass %.
4. A high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to claim 2, wherein the steel
sheet comprises C: 0.03-0.08 mass %, Si: 0.1-2.0 mass %, Mn:
1.0-3.0 mass %, P: not more than 0.05 mass % and S: not more than
0.01 mass %, provided that V, Nb and Ti satisfy a relationship of
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/48)- /(V/51).ltoreq.15.
5. A high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to any one of claims 1-4,
wherein the steel sheet further comprises one or two of the
following A-group and B-group: A-group: not more than 2.0 mass % in
total of one or two of Cr and Mo; B-group: not more than 2.0 mass %
in total of one or two of Cu and Ni.
6. A method of producing a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability, which comprises
hot rolling a steel slab having a composition comprising C:
0.01-0.08 mass %, Si: not more than 2.0 mass %, Mn: not more than
3.0 mass %, P: not more than 0.10 mass %, S: not more than 0.02
mass %, Al: 0.005-0.20 mass %, N: not more than 0.02 mass % and V:
0.01-0.5 mass %, provided that V and C satisfy a relationship of
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12, and the remainder
being Fe and inevitable impurities, pickling, cold rolling and then
subjecting to a continuous annealing at a temperature range from a
A.sub.C1 transformation point to a A.sub.C3 transformation
point.
7. A method of producing a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability, which comprises
hot rolling a steel slab having a composition comprising C:
0.01-0.08 mass %, Si: not more than 2.0 mass %, Mn: not more than
3.0 mass %, P: not more than 0.10 mass %, S: not more than 0.02
mass %, Al: 0.005-0.20 mass %, N: not more than 0.02 mass % and V:
0.01-0.5 mass % and further comprising not more than 0.3 mass % in
total of one or two of Nb: more than 0 mass % but not more than 0.3
mass % and Ti: more than 0 mass % but not more than 0.3 mass %,
provided that V, Nb, Ti and C satisfy a relationship of
0.5.times.C/12.ltoreq.(V/51+2.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C-
/12, and the remainder being Fe and inevitable impurities,
pickling, cold rolling and then subjecting to a continuous
annealing at a temperature range from a A.sub.C1 transformation
point to a A.sub.C3 transformation point.
8. A method of producing a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability according to claim
7, wherein the steel slab comprises not more than 0.3 mass % in
total of one or two of Nb: 0.001-0.3 mass % and Ti: 0.001-0.3 mass
%.
9. A method of producing a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability according to claim
7, wherein the steel slab comprises C: 0.03-0.08 mass %, Si:
0.1-2.0 mass %, Mn: 1.0-3.0 mass %, P: not more than 0.05 mass %
and S: not more than 0.01 mass %, provided that V, Nb and Ti
satisfy a relationship of
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/48)/(V/51).ltoreq.15.
10. A method of producing a high-strength dual-phase cold rolled
steel sheet having an excellent deep drawability according to any
one of claims 6-9, wherein the steel slab further comprises one or
two of the following A-group and B-group: A-group: not more than
2.0 mass % in total of one or two of Cr and Mo; B-group: not more
than 2.0 mass % in total of one or two of Cu and Ni.
11. A high-strength dual-phase galvanized steel sheet having an
excellent deep drawability comprising a galvanized coating on the
steel sheet as claimed in claim 1.
12. A high-strength dual-phase galvanized steel sheet having an
excellent deep drawability comprising a galvanized coating on the
steel sheet as claimed in claim 2.
13. A high-strength dual-phase galvanized steel sheet having an
excellent deep drawability comprising a galvanized coating on the
steel sheet as claimed in claim 3.
14. A high-strength dual-phase galvanized steel sheet having an
excellent deep drawability comprising a galvanized coating on the
steel sheet as claimed in claim 4.
15. A high-strength dual-phase galvanized steel sheet having an
excellent deep drawability comprising a galvanized coating on the
steel sheet as claimed in claim 5.
16. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability, characterized by
subjecting to a galvanization after the continuous annealing at a
temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point in the method claimed in claim 6.
17. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability according to claim
16, characterized by further comprising a continuous annealing step
between the cold rolling step and the continuous annealing step at
a temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point.
18. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability, characterized by
subjecting to a galvanization after the continuous annealing at a
temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point in the method claimed in claim 7.
19. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability according to claim
18, characterized by further comprising a continuous annealing step
between the cold rolling step and the continuous annealing step at
a temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point.
20. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability, characterized by
subjecting to a galvanization after the continuous annealing at a
temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point in the method claimed in claim 8.
21. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability according to claim
20, characterized by further comprising a continuous annealing step
between the cold rolling step and the continuous annealing step at
a temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point.
22. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability, characterized by
subjecting to a galvanization after the continuous annealing at a
temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point in the method claimed in claim 9.
23. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability according to claim
22, characterized by further comprising a continuous annealing step
between the cold rolling step and the continuous annealing step at
a temperature range from a A.sub.C1 transformation point to a
A.sub.C3 transformation point.
24. A method of producing a high-strength dual-phase galvanized
steel sheet having an excellent deep drawability according to any
one of claims 16-23, wherein the steel slab further comprises one
or two of the following A-group and B-group: A-group: not more than
2.0 mass % in total of one or two of Cr and Mo; B-group: not more
than 2.0 mass % in total of one or two of Cu and Ni.
Description
TECHNICAL FIELD
[0001] This invention relates to a high-strength dual-phase steel
sheet having an excellent deep drawability, and particularly to a
high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability and a high strength dual phase
galvanized steel sheet having an excellent deep drawability which
have a tensile strength of 440 MPa or more and are suitable for use
in steel sheets for vehicles as well as a method of producing the
same.
BACKGROUND ART
[0002] Recently, it is required to improve a fuel consumption in a
vehicle from a viewpoint of the maintenance of the global
environment, and also it is required to improve a safety of a
vehicle body from a viewpoint of the protection of crews during the
collision of the vehicle. To this end, investigations for achieving
both the lightening and strengthening of the vehicle body are
positively proceeding.
[0003] In order to simultaneously satisfy the lightening and
strengthening of the vehicle body, it is said that the
high-strengthening of raw materials constituting the parts is
effective, and recently, high-strength steel sheets are positively
used as a part of the vehicle.
[0004] Most of the parts for the vehicle body are formed by press
working of the steel sheet as a raw material. To this end, the
high-strength steel sheet used is required to have an excellent
press formability. In order to improve the press formability, it is
necessary to have a high Lankford value (r-value), a high ductility
(E1) and a low yield stress (YS) as mechanical properties of the
steel sheet.
[0005] However, in general, as the steel sheet becomes highly
strengthened, the r-value and the ductility lower and the press
formability is degraded, while the yield stress rises to degrade
the shapability and hence the problem of springback is apt to
occur.
[0006] And also, a high corrosion resistance is required according
to a position of the vehicle part to be applied, so that various
surface-treated steel sheets having an excellent corrosion
resistance are used as a steel sheet for the vehicle parts up to
now. Among these surface-treated steel sheets, a galvanized steel
sheet is manufactured in a continuous galvanizing equipment
conducting recrystallization annealing and galvanizing at the same
line, so that the provision of an excellent corrosion resistance
and a cheap production are possible. And also, an alloyed
galvanized steel sheet obtained by subjecting to a heat treatment
after the galvanization is excellent in the weldability and press
formability in addition to the excellent corrosion resistance.
Therefore, they are widely used.
[0007] In order to further advance the lightening and strengthening
of the vehicle body, in addition to the development of the
high-strength cold rolled steel sheet having the excellent press
formability, it is desired to develop a high-strength galvanized
steel sheet having an excellent corrosion resistance through the
continuous galvanizing line.
[0008] As a typical example of the high-strength steel sheet having
a good press formability is mentioned a dual-phase steel sheet
having a dual-phase microstructure of a soft ferrite phase and a
hard martensite phase. Especially, the dual-phase steel sheet
produced by cooling with a gas jet after the continuous annealing
is low in the yield stress and possesses a high ductility and an
excellent baking hardenability. The above dual-phase steel sheet is
generally good in the workability, but has a drawback that the
workability under severer condition is poor and particularly, the
r-value is low and the deep drawability is bad.
[0009] And also, when the galvanization is applied for providing
the excellent corrosion resistance, the continuous galvanizing line
is general to set up the annealing equipment and the plating
equipment continuously. To this end, in case of subjecting to the
galvanization, the cooling after the annealing is constrained by a
plating temperature and can not drop down to a temperature lower
than the plating temperature at once and hence the cooling is
interrupted. At a result, an average cooling rate necessarily
becomes smaller. Therefore, when the galvanized steel sheet is
produced in the continuous galvanizing line, it is difficult to
generate martensite phase produced under a cooling condition of a
large cooling rate into the steel sheet after the galvanization. To
this end, it is generally difficult to produce the high-strength
galvanized steel sheet having a dual-phase microstructure of a
ferrite phase and a martensite phase through the continuous
galvanizing line.
[0010] Under such unfavorable conditions, it is attempted to
increase the r-value of the dual-phase steel sheet to improve the
deep drawability. For example, JP-B-55-10650 discloses a technique
that a box annealing is carried out at a temperature ranging from a
recrystallization temperature to A.sub.c3 transformation point
after the cold rolling and thereafter the continuous annealing
inclusive of quenching and tempering is carried out after the
heating to 700-800.degree. C. in order to obtain the mixed
microstructure. In this method, however, the quenching and
tempering are carried out during the continuous annealing, so that
the yield stress is high and hence a low yield ratio can not be
obtained. The steel sheet having such a high yield stress is not
suitable for the press formability and has a drawback that the
shapability in the pressed parts is bad.
[0011] And also, a method for lowering the high yield stress is
disclosed in JP-A-55-100934. In this method, the box annealing is
first carried out in order to obtain a high r-value, wherein the
temperature in the box annealing is made to a two-phase region of
ferrite (.alpha.)-austenite (.gamma.) and Mn is enriched from
.alpha. phase to .gamma. phase during the soaking. As the Mn
enriched phase preferentially becomes .gamma. phase during the
continuous annealing, the dual-phase microstructure is obtained
even at a cooling rate as in the gas jet cooling, and further the
yield stress becomes low. In this method, however, it is required
to conduct the box annealing at a relatively high temperature being
the .alpha.-.gamma. two-phase region over a long time for enriching
Mn, so that there are many problems in production steps such as a
frequent occurrence of adhesion between steel sheets inside a coil
resulted from the thermal expansion in the annealing, an occurrence
of temper color, a lowering of service life in an inner cover for a
furnace body and the like. Therefore, it was difficult to
industrially stably produce high-strength steel sheets possessing a
high r-value and a low yield stress up to now.
[0012] In addition, JP-B-1-35900 discloses a technique wherein the
dual-phase cold rolled steel sheet having a very high r-value and a
low yield stress of r-value=1.61, YS=224 MPa and TS=482 MPa can be
produced by cold rolling a steel having a composition of 0.012 mass
% C-0.32 mass % Si-0.53 mass % Mn-0.03 mass % P-0.051 mass % Ti,
heating to 870.degree. C. corresponding to .alpha.-.gamma.
two-phase region and thereafter cooling at an average cooling rate
of 100.degree. C./s. However, the high cooling rate of 100.degree.
C./s is difficult to attain in the gas jet cooling usually used in
the continuous annealing line or continuous galvanizing line after
the cold rolling, and is required to use an equipment for
water-quenching, and also a problem becomes actual in the surface
treatment of the water-quenched steel sheet, so that there are
problems in the production equipment and the materials.
[0013] Furthermore, it is attempted to produce the high-strength
dual-phase galvanized steel sheet. In the past, as the method of
producing the high-strength dual-phase galvanized steel sheet is
generally used a method wherein the formation of low-temperature
transformation phase is facilitated by using a steel added with a
large amount of an alloying element such as Cr or Mo for enhancing
a hardenability. However, the addition of the large amount of the
alloying element undesirably brings about the rise of the
production cost.
[0014] Moreover, as is disclosed in JP-B-62-40405 and the like,
there is proposed a method of producing the high-strength
dual-phase galvanized steel sheet by defining the cooling rate
after the annealing or the plating in the continuous galvanizing
line. However, this method is not actual from the constraint on the
equipment for the continuous galvanizing line and also the steel
sheet obtained by this method is not said to have a sufficient
ductility.
DISCLOSURE OF THE INVENTION
[0015] It is, therefore, an object of the invention to solve the
aforementioned problems and to provide high-strength dual-phase
cold rolled steel sheets having an excellent deep drawability and
high-strength dual-phase galvanized steel sheets having an
excellent deep drawability as well as a method of producing the
same.
[0016] Moreover, the term "galvanized steel sheet" used herein
means to include a galvanized steel sheet obtained by subjecting to
a galvanization containing aluminum or the like in addition to zinc
and an alloyed galvanized steel sheet obtained by subjecting to a
heat (alloying) treatment for diffusing iron of the matrix steel
sheet into the plated layer after the galvanization.
[0017] In order to achieve the above object, the inventors have
made various studies with respect to an influence of the alloying
element upon the microstructure and the recrystallization texture
in the steel sheet. As a result, it has been found that by limiting
C in a steel slab to a lower content and rationalizing V content in
relation to C content, before the recrystallization annealing, C in
the steel is precipitated as a V carbide to decrease solid-solute C
as far as possible to thereby develop {111} recrystallization
texture to obtain a high r-value and subsequently the V carbide is
dissolved by heating to .alpha.-.gamma. two-phase region to enrich
C in austenite for easily generating martensite in a subsequent
cooling process, whereby the high-strength dual-phase cold rolled
steel sheet and high-strength dual-phase galvanized steel sheet
having a high r-value and an excellent deep drawability can be
produced stably.
[0018] The results of fundamental experiments performed by the
inventors will be explained below.
[0019] In this case, the experiments are performed with respect to
a high-strength dual-phase cold rolled steel sheet of TS: 590 MPa
grade and a high-strength dual-phase cold rolled steel sheet of TS:
780 MPa grade.
[0020] Firstly, the fundamental experiment in the high-strength
dual-phase cold rolled steel sheet of TS: 590 MPa grade is
performed under the following conditions. Each of various sheet
bars having a basic composition of C: 0.03 mass %, Si: 0.02 mass %,
Mn: 1.7 mass %, P: 0.01 mass %, S: 0.005 mass %, Al: 0.04 mass %
and N: 0.002 mass % and different V contents by adding V within a
range of 0.03-0.55 mass % is heated to 1250.degree. C. and soaked,
and then subjected to three-pass rolling at a finisher delivery
temperature of 900.degree. C. to obtain a hot rolled steel sheet
having a thickness of 4.0 mm.
[0021] In addition, the same procedure as described above is
conducted with respect to various sheet bars having a basic
composition of C: 0.03 mass %, Si: 0.02 mass %, Mn: 1.7 mass %, P:
0.01 mass %, S: 0.005 mass %, Al: 0.04 mass % and N: 0.002 mass %
and different values of (2.times.Nb [mass %]/93+2.times.Ti [mass
%]/48)/(V [mass %]/51) by adding V, Nb and Ti within ranges of
0.03-0.04 mass %, 0.01-0.18 mass % and 0.01-0.18 mass %,
respectively, so as to satisfy a relationship of 0.5.times.C [mass
%]/12.ltoreq.(V [mass %]/51+2.times.Nb [mass %]/93+2.times.Ti [mass
%]/48).ltoreq.3.times.C [mass %]/12.
[0022] Moreover, the hot rolled steel sheet after the finish
rolling is subjected to a temperature holding treatment of
650.degree. C..times.1 hour as a coiling treatment. Subsequently,
the sheet is subjected to a cold rolling at a rolling reduction of
70% to obtain a cold rolled steel sheet having a thickness of 1.2
mm. Next, the cold rolled steel sheet is subjected to a
recrystallization annealing at 850.degree. C. for 60 seconds and
cooled at a cooling rate of 30.degree. C./s.
[0023] On the other hand, the fundamental experiment in the
high-strength dual-phase cold rolled steel sheet of TS:780 MPa
grade is performed under the following conditions.
[0024] Each of various sheet bars having a basic composition of C:
0.04 mass %, Si: 0.70 mass %, Mn: 2.6 mass %, P: 0.04 mass %, S:
0.005 mass %, Al: 0.04 mass % and N: 0.002 mass % and different
values of (2.times.Nb/93+2.times.Ti/48)/(V/51) by adding V, Nb and
Ti within ranges of 0.02-0.06 mass %, 0.01-0.12 mass % and
0.01-0.12 mass %, respectively, so as to satisfy a relationship of
0.5.times.C [mass %]/12.ltoreq.(V [mass %]/51+2.times.Nb [mass
%]/93+2.times.Ti [mass %]/48).ltoreq.3.times.C [mass %]/12is heated
to 1250.degree. C. and soaked, and then subjected to three-pass
rolling at a finisher delivery temperature of 900.degree. C. to
obtain a hot rolled steel sheet having a thickness of 4.0 mm.
Moreover, the sheet after the finish rolling is subjected to a
temperature holding treatment of 650.degree. C..times.1 hour as a
coiling treatment. Subsequently, the sheet is subjected to a cold
rolling at a rolling reduction of 70% to obtain a cold rolled steel
sheet having a thickness of 1.2 mm. Next, the cold rolled steel
sheet is subjected to a recrystallization annealing at 850.degree.
C. for 60 seconds and cooled at a cooling rate of 30.degree.
C./s.
[0025] With respect to the thus obtained cold rolled steel sheets
is conducted out a tensile test to investigate tensile properties.
The tensile test is carried out by using JIS No. 5 tensile test
piece. The r-value is determined as an average r-value
{=(r.sub.L+.sub.C+2.times.r.s- ub.D)/4} in a rolling direction
(r.sub.L), a direction (r.sub.D) inclined at 45 degree with respect
to the rolling direction and a direction (r.sub.C) perpendicular
(90.degree.) to the rolling direction.
[0026] FIGS. 1a and 1b show an influence of V content in a steel
slab upon r-value and yield ratio of a cold rolled steel sheet
(YR=yield stress (YS)/tensile strength (TS).times.100(%)) in cold
rolled steel sheets of TS: 590 MPa grade produced by using a steel
slab containing V but not containing Nb and Ti, V. Moreover, an
abscissa in FIGS. 1a and 1b is an atomic ratio ((V/51)/(C/12)) of V
content to C content, and an ordinate is r-value in FIG. 1a and
yield ratio (YR) in FIG. 1b.
[0027] As seen from FIGS. 1a and 1b, a high r-value and a low yield
ratio are obtained by limiting V content in the steel slab to a
range of 0.5-3.0 as the atomic ratio to C content and it is
possible to produce high-strength dual-phase cold rolled steel
sheet having an excellent deep drawability.
[0028] In the steel sheet according to the invention, the inventors
found that a high r-value is obtained because solid-solute C and N
are less and {111} recrystallization texture is strongly developed
before the recrystallization annealing. And also, the inventors
found that by annealing at .alpha.-.gamma. two-phase region is
dissolved V carbide and the solid-solute C is enriched into
austenite phase in large quantity and the austenite can be easily
transformed into martensite in the subsequent cooling process to
obtain a dual-phase microstructure of ferrite and martensite.
[0029] Although Ti and Nb have mainly been used as a carbide
forming element in the past, the inventors paid notice to V having
a solubility of carbide higher than those of Ti and Nb for
effectively obtaining the solid-solute C by annealing at a higher
temperature region. That is, it is found that since V carbide
easily dissolves as compared with Ti carbide and Nb carbide in the
annealing at a high temperature, a sufficient amount of
solid-solute C for transforming austenite to martensite is obtained
by annealing at the .alpha.-.gamma. two-phase region. In addition,
it is clear that this phenomenon is most remarkably generated by V,
but the similar result is obtained by adding Nb and Ti
together.
[0030] Although the invention is based on the above knowledge, the
following knowledge is obtained to achieve another invention.
[0031] The inventors compared r-values in the high-strength
dual-phase cold rolled steel sheets of TS: 590 MPa grade and TS:
780 MPa produced by using steel slabs containing Nb and Ti in
addition to V and made clear the followings. FIGS. 2a and 2b show
an influence of V, Nb and Ti contents in the steel slab upon
tensile strength (TS) and Lankford value (r-value) of a cold rolled
steel sheet in the cold rolled steel sheets of TS: 590 MPa grade
and TS: 780 MPa grade produced by using the V, Nb and Ti containing
steel slab. Moreover, an abscissa in FIGS. 2a and 2b is an atomic
ratio (2.times.Nb/93+2.times.Ti/48)/(V/51) of Nb and Ti contents to
V content, and an ordinate is tensile strength (TS) in FIG. 2a and
r-value in FIG. 2b.
[0032] According to the above results, in the TS: 780 MPa grade,
the high-strengthening is attempted by large quantities of
solid-solution strengthening elements, so that the r-value is
lowered as compared with that of the TS: 590 MPa grade by the
increase of the solid-solute C content or the like. In the TS: 780
MPa grade, however, the r-value is considerably improved when the
value of (2.times.Nb/93+2.times.Ti/48)/(V/- 51) is a range of not
less than 1.5. Such a characteristic in the TS: 780 MPa grade that
the r-value is remarkably improved when the value of
(2.times.Nb/93+2.times.Ti/48)/(V/51) is a range of not less than
1.5 is not recognized in the TS: 590 MPa grade.
[0033] Although the detail of causes on the above result is not
clear, it is considered that in the system containing a large
amount of an element resulted in the lowering of the r-value such
as solid-solute C or the like as in the TS: 780 MPa grade, Nb and
Ti easily precipitate the solid-solute C and N as a compound as
compared with V and the solid-solute C and N contents after the hot
rolling become less to improve the r-value. Moreover, when the
value of (2.times.Nb/93+2.times.T- i/48)/(V/51) exceeds 15, TS
considerably lowers, which is unfavorable for obtaining the
high-strength dual-phase cold rolled steel sheet of TS: 780 MPa
grade. This is considered due to the fact that as Nb carbide and Ti
are hardly dissolved as compared with V carbide, if the addition
quantities of the Nb and Ti contents are larger than that of the V
content, the C content enriched in austenite phase is largely
decreased in the annealing at the .alpha.-.gamma. two-phase region
is widely decreased and martensite phase generated after the
cooling is softened.
[0034] The invention is accomplished by further examining based on
the above knowledge. The summary of the invention is as
follows.
[0035] (1) A high-strength dual-phase cold rolled steel sheet
having an excellent deep drawability, characterized in that the
steel sheet has a composition comprising C: 0.01-0.08 mass %, Si:
not more than 2.0 mass %, Mn: not more than 3.0 mass %, P: not more
than 0.10 mass %, S: not more than 0.02 mass %, Al: 0.005-0.20 mass
%, N: not more than 0.02 mass % and V: 0.01-0.5 mass % provided
that V and C satisfy a relationship represented by the following
equation (i):
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12 (i)
[0036] and the remainder being Fe and inevitable impurities, and
has a microstructure consisting of a ferrite phase as a primary
phase and a secondary phase including martensite phase at an area
ratio of not less than 1% to a whole of the microstructure.
[0037] (2) A high-strength dual-phase cold rolled steel sheet
having an excellent deep drawability according to the item (1),
wherein the steel sheet has a composition comprising further not
more than 0.3 mass % in total of one or tow of Nb: more than 0 mass
% but not more than 0.3 mass % and Ti: more than 0 mass % but not
more than 0.3 mass % provided that V, Nb, Ti and C satisfy a
relationship represented by the following equation (ii) instead of
the equation (i):
0.5.times.C/12.ltoreq.(V/51+2.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C/-
12 (ii)
[0038] and the remainder being Fe and inevitable impurities.
[0039] Moreover, it is preferable that one or two of Nb: 0.001-3.0
mass % and Ti: 0.001-0.3 mass % is not more than 0.3 mass % in
total.
[0040] (3) A high-strength dual-phase cold rolled steel sheet
having an excellent deep drawability according to the item (2),
wherein the steel sheet comprises C: 0.03-0.08 mass %, Si: 0.1-2.0
mass %, Mn: 1.0-3.0 mass %, P: not more than 0.05 mass % and S: not
more than 0.01 mass % and V, Nb and Ti satisfy a relationship of
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/- 48)/(V/51) .ltoreq.15.
[0041] (4) A high-strength dual-phase cold rolled steel sheet
having an excellent deep drawability according to any one of the
items (1) to (3), wherein the steel sheet further comprises one or
two of the following A group and B group:
[0042] A group: not more than 2.0 mass % in total of one or two of
Cr and Mo;
[0043] B group: not more than 2.0 mass % in total of one or two of
Cu and Ni.
[0044] (5) A method of producing a high-strength dual-phase cold
rolled steel sheet having an excellent deep drawability, which
comprises hot rolling a steel slab having a composition comprising
C: 0.01-0.08 mass %, Si: not more than 2.0 mass %, Mn: not more
than 3.0 mass %, P: not more than 0.10 mass %, S: not more than
0.02 mass %, Al: 0.005-0.20 mass %, N: not more than 0.02 mass %
and V: 0.01-0.5 mass % provided that V and C satisfy a relationship
represented by the following equation (iii):
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12 (iii)
[0045] and the remainder being Fe and inevitable impurities,
pickling, cold rolling and then subjecting to a continuous
annealing at a temperature range from a A.sub.C1 transformation
point to a A.sub.C3 transformation point.
[0046] (6) A method of producing a high-strength dual-phase cold
rolled steel sheet having an excellent deep drawability according
to the item (5), wherein the steel sheet has a composition
comprising further not more than 0.3 mass % in total of one or tow
of Nb: more than 0 mass % but not more than 0.3 mass % and Ti: more
than 0 mass % but not more than 0.3 mass % provided that V, Nb, Ti
and C satisfy a relationship represented by the following equation
(iv) instead of the equation (iii):
0.5.times.C/12.ltoreq.(V/51+2.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C/-
12 (iv)
[0047] and the remainder being Fe and inevitable impurities.
[0048] Moreover, it is preferable that one or two of Nb: 0.001-0.3
mass % and Ti: 0.001-0.3 mass % is not more than 0.3 mass % in
total.
[0049] (7) A method of producing a high-strength dual-phase cold
rolled steel sheet having an excellent deep drawability according
to the item (6), wherein the steel slab comprises C: 0.03-0.08 mass
%, Si: 0.1-2.0 mass %, Mn: 1.0-3.0 mass %, P: not more than 0.05
mass % and S: not more than 0.01 mass % and V, Nb and Ti satisfy a
relationship of
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/48)/(V/51).ltoreq.15.
[0050] (8) A method of producing a high-strength dual-phase cold
rolled steel sheet having an excellent deep drawability according
to any one of the items (5)-(7), wherein the steel slab further
comprises one or two of the following A-group and B-group:
[0051] A-group: not more than 2.0 mass % in total of one or two of
Cr and Mo;
[0052] B-group: not more than 2.0 mass % in total of one or two of
Cu and Ni.
[0053] (9) A high-strength dual-phase galvanized steel sheet having
an excellent deep drawability comprising a galvanized coating on
the steel sheet disclosed in any one of the items (1)-(4).
[0054] (10) A method of producing a high-strength dual-phase
galvanized steel sheet having an excellent deep drawability,
wherein a galvanization is carried out after the continuous
annealing at a temperature range from a A.sub.C1 transformation
point to a A.sub.C3 transformation point in the production method
described in any one of the items (5)-(7).
[0055] (11) A method of producing a high-strength dual-phase
galvanized steel sheet having an excellent deep drawability
according to the item (10), which further comprising a continuous
annealing step between the cold rolling step and the continuous
annealing step at a temperature range from a A.sub.C1
transformation point to a A.sub.C3 transformation point.
[0056] (12) A method of producing a high-strength dual-phase
galvanized steel sheet having an excellent deep drawability
according to the item (10) or (11), wherein the steel slab further
comprises one or two of the following A-group and B-group:
[0057] A-group: not more than 2.0 mass % in total of one or two of
Cr and Mo;
[0058] B-group: not more than 2.0 mass % in total of one or two of
Cu and Ni.
[0059] The cold rolled steel sheet and the galvanized steel sheet
according to the invention are high-strength dual-phase steel
sheets having a tensile strength (TS) of not less than 440 MPa and
an excellent deep drawability.
[0060] At first, the reason of limiting the composition in the cold
rolled steel sheet and the galvanized steel sheet according to the
invention will be explained below. Moreover, mass % represents
simply as "%".
[0061] C: 0.01-0.08%
[0062] C is an element for increasing the strength of the steel
sheet and further promoting the formation of a dual-phase
microstructure of ferrite and martensite, and is necessary to
contain not less than 0.01%, preferably not less than 0.015% from a
viewpoint of the formation of the dual-phase microstructure in the
invention. Moreover, if it is intended to increase the strength to
TS: not less than 540 MPa and TS: not less than 780 MPa, the C
content is preferable to be not less than 0.015% and not less than
0.03%, respectively. On the other hand, when the C content exceeds
0.08%, the development of {111} recrystallization texture is
obstructed to degrade the deep drawability. Therefore, the
invention limits the C content to 0.01-0.08%. When it is
particularly required to increase the strength of the steel sheet,
it is preferable to be 0.03-0.08%. Moreover, it is preferable to be
not more than 0.05% from a viewpoint of the deep drawability.
[0063] Si: Not More Than 2.0%
[0064] Although Si is a useful reinforcing element capable of
increasing the strength of the steel sheet without remarkably
lowering the ductility of the steel sheet, if the content exceeds
2.0%, the deterioration of the deep drawability is caused, but also
the surface properties are degraded. Therefore, Si is limited to
not more than 2.0%. Moreover, if it is intended to increase the
strength to TS: not less than 780 MPa, it is preferable to be not
less than 0.1% for ensuring the required strength. And also, it is
preferable to be not less than 0.01% for increasing the strength to
TS: not less than 440 MPa which is a main object of the
invention.
[0065] Mn: Not More Than 3.0%
[0066] Mn has an action reinforcing the steel and further has an
action of lessening a critical cooling rate for the obtention of
the dual-phase microstructure of ferrite and martensite to promote
the formation of the dual-phase microstructure of ferrite and
martensite, so that it is preferable to contain a content in
accordance with the cooling rate after the recrystallization
annealing. And also, Mn is an effective element preventing the hot
tearing through S, so that it is preferable to contain an
appropriate content in accordance with S content. However, when the
Mn content exceeds 3.0%, the deep drawability and weldability are
degraded. In the invention, therefore, the Mn content is limited to
not more than 3.0%. Moreover, the Mn content is preferable to be
not less than 0.5% for remarkably developing the above effect, and
particularly it is preferable to be not less than 1.0% for
increasing the strength to TS: not less than 780 MPa. And also, it
is preferable to be not less than 0.1% for increasing the strength
to TS: not less than 440 MPa which is a main object of the
invention.
[0067] P: Not More Than 0.10%
[0068] P has an action reinforcing the steel and can be contained
in a required amount in accordance with the desired strength. When
the P content exceeds 0.10%, the press formability is degraded.
Therefore, the P content is limited to not more than 0.10%.
Moreover, if a more excellent press formability is required, the P
content is preferable to be not more than 0.08%. Furthermore, when
large quantities of C, Mn and the like are contained in order to
ensure TS: not less than 780 MPa, the P content is preferable to be
not more than 0.05% in order to prevent the degradation of the
weldability. In addition, if it is intended to increase the
strength to TS: not less than 440 MPa, it is preferable to be not
less than 0.001%.
[0069] S: Not More Than 0.02%
[0070] S is existent as an inclusion in the steel sheet and is an
element bringing about the degradation of the ductility and the
formability of the steel sheet, particularly the stretch-flanging
property. Therefore, it is preferable to be decreased as far as
possible, and when it is decreased to not more than 0.02%, S does
not exert a bad influence, so that the S content is 0.02% as an
upper limit in the invention. Moreover, when the more excellent
stretch-flanging property is required, or when the large quantities
of C, Mn and the like are contained in order to ensure TS: not less
than 780 MPa, if the excellent weldability is required, the S
content is preferable to be not more than 0.01%, more preferably
not more than 0.005%. On the other hand, the S content is
preferable to be not less than 0.0001% considering a cost for the
removal of S in the steelmaking process.
[0071] Al: 0.005-0.20%
[0072] Al is added to the steel as a deoxidizing element and is a
useful element for improving the cleanliness of the steel, but the
addition effect is not obtained at less than 0.005%. On the other
hand, when it exceeds 0.20%, the more deoxidizing effect is not
obtained and the deep drawability is inversely degraded. Therefore,
the Al content is limited to 0.005-0.20%. Moreover, the invention
does not exclude a steelmaking method through deoxidization other
than the Al deoxidization. For example, Ti deoxidization or Si
deoxidization may be conducted. The steel sheets made by these
deoxidizing methods are included within a scope of the invention.
In this case, even if Ca, REM and the like are added to the molten
steel, the characteristics of the steel sheet according to the
invention are not obstructed, so that the steel sheet including Ca,
REM and the like is naturally included within the scope of the
invention.
[0073] N: Not More Than 0.02%
[0074] N is an element increasing the strength of the steel sheet
by the solid-solution hardening and the strain ageing hardening,
but when N content exceeds 0.02%, the nitride is increased in the
steel sheet to remarkably degrade the deep drawability of the steel
sheet. Therefore, the N content is limited to not more than 0.02%.
Moreover, in case of requiring the more improvement of the press
formability, the N content is preferable to be not more than 0.01%,
more preferably not more than 0.004%. In this case, considering the
cost for denitrification in the steelmaking process, the N content
is preferable to be not less than 0.0001%.
[0075] V: 0.01-0.5% and
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12
[0076] V is a most important element in the invention. Before the
recrystallization, the solid-solute C is precipitated and fixed as
V carbide to develop the {111} recrystallization texture, whereby a
high r-value can be obtained. Moreover, V dissolves the V carbide
in the annealing at .alpha.-.gamma. two-phase region to enrich a
large quantity of the solid-solute C in austenite phase, which is
easily transformed into martensite at the subsequent cooling
process, whereby the dual-phase steel sheet having a dual-phase
microstructure of ferrite and martensite can be obtained. Such an
effect becomes effective when the V content is not less than 0.01%,
more preferably not less than 0.02% and satisfies
0.5.times.C/12.ltoreq.V/51 in relation to the C content. On the
other hand, when the V content exceeds 0.5% or when it is
V/51>3.times.C/12 in relation to the C content, the dissolution
of the V carbide at the .alpha.-.gamma. two-phase region hardly
occurs and the dual-phase microstructure of ferrite and martensite
is hardly obtained. Therefore, the V content is limited to
0.01-0.5% and to 0.5.times.C/12.ltoreq.V/51.l- toreq.3.times.C/12.
Moreover, V/51.ltoreq.2.times.C/12 is preferable for obtaining the
dual-phase microstructure of ferrite and martensite.
[0077] In addition to the above composition, it is further
preferable to contain not more than 0.3 (mass) % in total of one or
two of Nb: more than 0% but not more than 0.3 (mass) % and Ti: more
than 0% but not more than 0.3%, and that V, Nb, Ti contents satisfy
0.5.times.C/12.ltoreq.(V/5-
1+2.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C/12 in relation to
the C content in place of that the V and C content satisfy
0.5.times.C/12.ltoreq.V/51.ltoreq.3.times.C/12.
[0078] Not More Than 0.3% in Total of One or Tow of Nb: More Than
0% but not More Than 0.3% and Ti: More Than 0% but not More Than
0.3%, and V, Nb, Ti and C Satisfy
0.5.times.C/12.ltoreq.(V/51+2-Nb/93+2.times.Ti/48).l-
toreq.3.times.C/12
[0079] Nb and Ti are carbide forming elements likewise V and have
the same action as V mentioned above. That is, a high r-value can
be obtained by precipitating and fixing the solid-solute C as Nb
and Ti carbides before the recrystallization to develop the {111}
recrystallization texture, and also a dual-phase steel sheet having
a dual-phase microstructure of ferrite and martensite can be
obtained by dissolving the Nb and Ti carbides in the annealing at
the .alpha.-.gamma. two-phase region to enrich a large quantity of
the solid-solute C in austenite phase and transforming into
martensite in the subsequent cooling process. Moreover, as the
above effect of Nb and Ti is considerably small as compared with
that of V, when only Nb and Ti are added to the steel slab without
adding V, the deep drawability aiming at the invention can not be
enhanced sufficiently.
[0080] Therefore, it is preferable to add Nb and Ti of more than
0%. More preferably, each of the Nb and Ti contents is not less
than 0.001%. In this case, it is preferable to satisfy
0.5.times.C/12.ltoreq.(V/51+2.time- s.Nb/93+2.times.Ti/48) in
relation to the C and V contents for developing the above effect.
On the other hand, when each of Nb and Ti contents or both in total
thereof exceeds 0.3%, or when the Nb and Ti contents satisfy
(V/51+2.times.Nb/93+2.times.Ti/48)>3.times.C/12 in relation to
the C and V contents, the dissolution of the carbide at the
.alpha.-.gamma. two-phase region hardly occurs and hence the
dual-phase microstructure of ferrite and martensite is hardly
obtained. Therefore, it is preferable that when either Nb or Ti is
merely added, each of the Nb content and the Ti content is within a
range of more than 0% but not more than 0.3%, and when both of Nb
and Ti are added together, the Nb and Ti contents are not more than
0.3% in total and satisfy
0.5.times.C/12.ltoreq.(V/51+2.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C-
/12 in relation to the V and C contents.
[0081] On the other hand, if it is intended to increase the
strength to TS: not less than 780 MPa, the deep drawability is apt
to be easily degraded by the addition of large quantities of
solid-solution strengthening elements such as C, Mn and the like.
In this case, the V, Nb and Ti contents are further desirable to be
a range of
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/48)/(V/51).ltoreq.15. The
reason why (2.times.Nb/93+2.times.Ti/48)/(V/51) is limited to not
less than 1.5 is considered due to the fact that although the
detail of the cause is not clear, the formation of carbide after
the hot rolling is promoted to decrease the solid-solute C by
adding large quantities of Nb and Ti as compared with V and hence
the {111} recrystallization texture is easily developed. Moreover,
in order to ensure the strength of TS: not less than 780 MPa,
(2.times.Nb/93+2.times.Ti/48)/(V/51) is desirable to be a range of
not more than 15.
[0082] Furthermore, in addition to the above steel composition, the
steel according to the invention is preferable to further comprise
one or two of the following A-group and B-group:
[0083] A-group: not more than 2.0% in total of one or two of Cr and
Mo;
[0084] B-group: not more than 2.0% in total of one or two of Cu and
Ni.
[0085] A-Group: Not More Than 2.0% in Total of One or Two of Cr and
Mo
[0086] All of Cr and Mo in the A-group have an action of decreasing
the critical cooling rate for providing the dual-phase
microstructure of ferrite and martensite to promote the formation
of the dual-phase microstructure of ferrite and martensite likewise
Mn and can be included, if necessary. The lower limits of the Cr
content and Mo content preferable for obtaining the above effect
are Cr: 0.05%, Mn: 0.05%. However, when one or two of Cr and Mo
exceed 2.0% in total, the deep drawability is degraded. To this
end, one or more of Cr and Mo in the A-group is preferable to be
limited to not more than 2.0% in total.
[0087] B-Group: Not More Than 2.0% in Total of One or Two of Cu and
Ni
[0088] Cu and Ni in the B-group have an action of reinforcing the
steel and may be included at a required amount in accordance with
the desired strength. However, when the content of Cu and Ni added
alone or together exceeds 2.0% in total, it tends to degrade the
deep drawability. To this end, one or more of Cu and Ni is
preferable to be not more than 2.0% in total. Moreover, the lower
limits of the Cu and Ni contents preferable for obtaining the above
effect is Cu: 0.05% and Ni: 0.05%, respectively.
[0089] Although elements other than the above elements are not
particularly limited in the invention, there is no problem even if
B, Ca, Zr, REM and the like is included within a range of the usual
steel composition.
[0090] In this case, B is an element having an action of improving
the hardenability in the steel and may be included, if necessary.
However, when the B content exceeds 0.003%, the above effect is
saturated, so that the B content is preferable to be not more than
0.003%. Moreover, a more desirable range is 0.001-0.002%. Ca and
REM have an action of controlling the form of sulfide inclusion and
also have an effect of improving the stretch-flanging property.
Such an effect is saturated when one or two selected from Ca and
REM exceed 0.01% in total. To this end, the content of one or two
of Ca and REM is preferable to be not more than 0.01% in total.
Moreover, a more preferable range is 0.001-0.005%.
[0091] The reminder other than the above elements is Fe and
inevitable impurities. As the inevitable impurity are mentioned,
for example, Sb, Sn, Zn, Co and the like. As acceptable ranges of
their contents are 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%.
[0092] Next, the microstructure of the steel sheet according to the
invention will be explained.
[0093] The cold rolled steel sheet according to the invention has a
microstructure consisting of ferrite phase as a primary phase and a
secondary phase including not less than 1% of martensite phase at
an area ratio with respect to a whole of the microstructure.
[0094] In order to provide the cold rolled steel sheet having a low
yield stress (YS), a high ductility (E1) and an excellent deep
drawability, it is required to render the microstructure of the
steel sheet according to the invention into a dual-phase
microstructure consisting of a ferrite phase as a primary phase and
a secondary phase including a martensite phase. It is preferable
that the ferrite phase as a primary phase is not less than 80% at
an area ratio and hence the secondary phase is not more than 20%.
When the area ratio of the ferrite phase is less than 80%, it is
difficult to ensure the high ductility and the press formability
tends to lower. And also, when a good ductility is required, it is
preferable that the ferrite phase is not less than 85% at the area
ratio and hence the secondary phase is not more than 15%. Moreover,
in order to utilize the advantage of the dual-phase microstructure,
the ferrite phase is required to be not more than 99%.
[0095] In the invention, the secondary phase is required to include
the martensite phase at the area ratio of not less than 1% with
respect to the whole of the microstructure. When the martensite is
less than 1% at the area ratio, the low yield stress (YS) and the
high ductility (E1) can not be satisfied simultaneously. More
preferably, the martensite phase is not less than 3% but not more
than 20% at the area ratio. In case of requiring a good ductility,
the martensite phase is preferable to be not more than 15% at the
area ratio. Moreover, the secondary phase may be constituted by
only the martensite phase at the area ratio of not less than 1% or
by mixed phases of the martensite phase at the area ratio of not
less than 1% and any of a pearlite phase, a bainite phase and a
retained austenite as an additional phase and is not especially
limited. In the latter case, the pearlite phase, the bainite phase
and the retained austenite are preferable to be not more than 50%
in total at the area ratio with respect to the microstructure of
the secondary phase in order to more effectively develop the effect
of the martensite phase.
[0096] The cold rolled steel sheet and the galvanized steel sheet
having the above microstructure are steel sheets having a low yield
stress, a high ductility and an excellent deep drawability.
[0097] Next, the method of producing the cold rolled steel sheet
and the galvanized steel sheet according to the invention will be
explained.
[0098] The composition of the steel slab used in the production
method of the invention is the same as the compositions of the
aforementioned cold rolled steel sheet and the galvanized steel
sheet, so that the explanation on the reason of the limitation in
the steel slab is omitted.
[0099] The cold rolled steel sheet according to the invention is
produced by using a steel slab having a composition of the above
range as a starting material and successively subjecting this
starting material to a hot rolling step of subjecting to a hot
rolling to obtain a hot rolled steel sheet, a pickling step of
pickling the hot rolled steel sheet, a cold rolling step of
subjecting the hot rolled steel sheet to a cold rolling to obtain a
cold rolled steel sheet, and a recrystallization annealing step of
subjecting the cold rolled steel sheet to a recrystallization
annealing to obtain a cold rolled annealed steel sheet.
[0100] And also, the galvanized steel sheet according to the
invention is produced by using a steel slab having a composition of
the above range as a starting material and successively subjecting
this starting material to a hot rolling step of subjecting to a hot
rolling to obtain a hot rolled steel sheet, a pickling step of
pickling the hot rolled steel sheet, a cold rolling step of
subjecting the hot rolled steel sheet to a cold rolling to obtain a
cold rolled steel sheet, and a continuous galvanization step of
subjecting the cold rolled steel sheet to a recrystallization
annealing and a galvanizing to obtain a galvanized steel sheet.
Furthermore, it is produced by subjecting the cold rolled steel
sheet to a step of annealing and pickling before the continuous
galvanization step, if necessary.
[0101] The steel slab used is preferable to be produced by a
continuous casting process in order to prevent the
macro-segregation of the components, but may be produced by an
ingot casting process or a thin slab casting process. Furthermore,
in addition to the conventional process of cooling to a room
temperature once after the production of the steel slab and again
heating, energy-saving processes such as a process for inserting a
hot steel slab into a heating furnace without cooling, a process
for direct sending rolling or direct rolling immediately after
slight heat-holding and the like can be applied without
problems.
[0102] The above starting material (steel slab) is subjected to the
hot rolling step of forming the hot rolled steel sheet by heating
and hot rolling. In the hot rolling step, there is particularly no
problem even in the use of usual rolling conditions as long as the
hot rolled steel sheet having a desired thickness can be produced.
Moreover, preferable hot rolling conditions are mentioned below for
the reference.
[0103] Slab Heating Temperature: not Lower Than 900.degree. C.
[0104] The slab heating temperature is desirable to be made lower
as far as possible in order to improve the deep drawability by
coarsening the precipitate to develop the {111} recrystallization
texture. However, when the slab heating temperature is lower than
900.degree. C., the rolling load increases and the risk of causing
troubles in the hot rolling increases. To this end, the slab
heating temperature is preferable to be not lower than 900.degree.
C. And also, the upper limit of the slab heating temperature is
more preferable to be 1300.degree. C. in terms of the lowering of
the yield resulted from the increase of scale loss accompanied with
the increase of the oxide weight. Moreover, it goes without saying
that the utilization of a so-called sheet bar heater of heating the
sheet bar in the hot rolling is an effective process from a
viewpoint that the slab heating temperature is lowered and the
troubles in the hot rolling are prevented.
[0105] Finisher Delivery Temperature: not Lower Than 700.degree.
C.
[0106] The finisher delivery temperature (FDT) is preferable to be
not lower than 700.degree. C. in order to obtain a uniform
microstructure of the hot rolled parent sheet for providing an
excellent deep drawability after the cold rolling and the
recrystallization annealing. That is, when the finish deformation
temperature is lower than 700.degree. C., not only the
microstructure of the hot rolled parent sheet becomes nonuniform,
but also the rolling load in the hot rolling becomes higher and the
risk of causing the trouble in the hot rolling is increased.
[0107] Coiling Temperature: not More Than 800.degree. C.
[0108] The coiling temperature is preferable to be not higher than
800.degree. C. That is, when the coiling temperature exceeds
800.degree. C., the scale increases and the yield tends to lower
due to the scale loss. And also, when the coiling temperature is
lower than 200.degree. C., the shape of the steel sheet remarkably
is disordered and the risk of causing problems in the actual use
increases, so that the lower limit of the coiling temperature is
more preferable to be 200.degree. C.
[0109] As mentioned above, in the hot rolling step according to the
invention, it is preferable that the steel slab is heated above
900.degree. C., subjected to the hot rolling at the finish
deformation temperature of not lower than 700.degree. C., and
coiled at the coiling temperature of not higher than 800.degree.
C.
[0110] Moreover, in the hot rolling step according to the
invention, a lubrication rolling may be conducted in a part of the
finish rolling or between passes thereof in order to reduce the
rolling load in the hot rolling. In addition, the application of
the lubrication rolling is effective from a viewpoint of the
uniformization of the steel sheet shape and the homogenization of
the material. Also, the coefficient of friction in the lubrication
rolling is preferable to be within a range of 0.10-0.25.
[0111] Further, the hot rolling step is preferable to be a
continuous rolling process wherein the sheet bars located in front
and rear are joined to each other and continuously subjected to the
finish rolling. The application of the continuous rolling process
is desirable from a viewpoint of the operating stability in the hot
rolling.
[0112] Next, the hot rolled steel sheet is subjected to the
pickling for the removal of the scale. The pickling step is
sufficient according to the usual manner and it is preferable to
use a treating solution such as hydrochloric acid, sulfuric acid or
the like as a pickling solution.
[0113] Moreover, the cold rolled steel sheet is formed by
subjecting the hot rolled steel sheet to the cold rolling. The cold
rolling conditions are not especially limited as long as the cold
rolled steel sheet having desired size and shape can be obtained,
but it is preferable that a rolling reduction in the cold rolling
is not less than 40%. When the rolling reduction is less than 40%,
the {111} recrystallization texture is not developed and the
excellent deep drawability can not be obtained.
[0114] The cold rolled steel sheet according to the invention is
subjected to a recrystallization annealing in the subsequent
recrystallization annealing step to obtain a cold rolled annealed
steel sheet. The recrystallization annealing is carried out in a
continuous annealing line. On the other hand, the galvanized steel
sheet according to the invention is produced by subjecting the cold
rolled steel sheet to recrystallization annealing and galvanizing
in the continuous galvanization line after the cold rolling. In
this case, the annealing temperature in the recrystallization
annealing is required to be conducted at a (.alpha.+.gamma.)
two-phase region within a temperature range from A.sub.C1
transformation point to A.sub.C3 transformation point. This is due
to the fact that the annealing is carried out at (.alpha.+.gamma.)
two-phase region to dissolve the carbides of V, Ti and Nb to
thereby distribute an amount of solid-solute C sufficient to
transform austenite to martensite into the austenite phase. When
the annealing temperature is lower than the A.sub.C1 transformation
point, the microstructure is rendered into the ferrite single phase
and the martensite can not be generated, while when it is higher
than the A.sub.C3 transformation point, the crystal grains are
coarsened and the microstructure is rendered into the austenite
single phase and the {111} recrystallization texture is not
developed and hence the deep drawability is deteriorated
remarkably.
[0115] In the cold rolled steel sheet according to the invention,
the cooling in the recrystallization annealing is preferable to be
conducted at a cooling rate of not less than 5.degree. C./s in
order to produce the martensite phase to obtain the dual-phase
microstructure of ferrite and martensite.
[0116] On the other hand, in the galvanized steel sheet according
to the invention, it is preferable to quench to a temperature
region of 380-530.degree. C. after the above recrystallization
annealing. When a stop temperature of the quenching is lower than
380.degree. C., the defective plating easily occurs, while when it
exceeds 530.degree. C., the unevenness easily occurs on the plated
surface. Moreover, the cooling rate is preferable to be not less
than 5.degree. C./s in order to produce the martensite phase to
obtain the dual-phase microstructure of ferrite and martensite.
After the above quenching, the galvanization is carried out by
dipping in a galvanizing bath. In this case, Al concentration in
the galvanizing bath is preferable to be within a range of
0.12-0.145 mass %. When the Al concentration in the galvanizing
bath is less than 0.12 mass %, the alloying excessively advances
and the plating adhesion (resistance to powdering) tends to be
deteriorated, while when it exceeds 0.145 mass %, the defective
plating easily occurs.
[0117] And also, the plated layer may be subjected to an alloying
treatment after the galvanization. Moreover, the alloying treatment
is preferable to be conducted so that Fe content in the plated
layer is 9-12%.
[0118] As the alloying treatment, it is preferable to conduct the
alloying of the galvanized layer by reheating up to a temperature
region of 450-550.degree. C. After the alloying treatment, it is
preferable to cool at a cooling rate of not less than 5.degree.
C./s to 300.degree. C. The alloying at a high temperature is
difficult to form the martensite phase and there is caused a fear
of degrading the ductility of the steel sheet, while when the
alloying temperature is lower than 450.degree. C., the progress of
the alloying is slow and the productivity tends to lower.
Furthermore, when the cooling rate after the alloying treatment is
extremely small, the formation of the martensite becomes difficult.
To this end, the cooling rate at a temperature region from after
the alloying treatment to 300.degree. C. is preferable to be not
less than 5.degree. C./s.
[0119] Moreover, if it is required to further improve the plating
property, it is preferable that after the cold rolling and before
being subjected to the continuous galvanization, the annealing is
separately conducted in the continuous annealing line and
subsequently an enriched layer of components in the steel produced
on the surface of the steel sheet is removed by pickling and
thereafter the above treatment is conducted in the continuous
galvanization line. In this case, the pickling may be carried out
in the pickling line or in the pickling bath arranged in the
continuous galvanization line. Also, the atmosphere in the
continuous annealing line is preferable to be a reducing atmosphere
with respect to the steel sheet in order to prevent the formation
of the scale, and it is generally sufficient to use a nitrogen gas
containing several % of H.sub.2. The annealing is preferable to be
conducted under a condition that a temperature of the steel sheet
reaching in the continuous annealing line is not lower than the
A.sub.C1 transformation point decided by the components in the
steel. Because it is required to promote the enrichment of the
alloying element on the surface of the steel sheet and to enrich
the alloying element in the secondary phase by once forming the
dual-phase microstructure in the continuous annealing line. In the
steel sheet after the annealing in the continuous annealing line,
there is a tendency that P among the components in the steel is
diffused to segregate on the surface of the steel sheet and Si, Mn,
Cr and the like enrich as an oxide, so that it is preferable to
remove the enriched layer formed on the surface of the steel sheet
by the pickling. Then, the same annealing as in the above is
performed in the continuous galvanization line. In order to develop
the characteristics as the dual-phase microstructure, the annealing
in the continuous galvanization line is preferable to be performed
at (.alpha.+.gamma.) two-phase region within a temperature range of
from the A.sub.C1 transformation point to the A.sub.C3
transformation point. In this case, the reason why the annealing is
performed at not lower than the A.sub.C1 transformation point in
both the continuous annealing line and the continuous galvanization
line is due to the fact that the dual-phase microstructure is
formed as mentioned above. Once an enriching place of the element
as the secondary phase is formed by forming the dual-phase
microstructure as a final microstructure in the continuous
annealing line, it becomes possible to enrich the alloying element
to some degree at this place. Desirably, it is sufficient to obtain
the same dual-phase microstructure as in a final product after the
cooling, so that the alloying element is more preferable to be
enriched in the vicinity of a triple point of grain boundary
(intersection of the grain boundary formed by three crystal
grains). Thereafter, when the annealing is performed at the
two-phase region in the continuous galvanization line, the alloying
element is further enriched in the secondary phase or .gamma.-phase
and hence the .gamma.-phase easily transforms into the martensite
phase during the cooling process. Moreover, the term "alloying
element" used herein means a substitutional alloying element such
as Mn, Mo or the like, which makes a situation that diffusion
hardly occurs and enrichment easily occurs at the temperature in
the annealing step in order to lower the yield ratio.
[0120] And also, the cold rolled steel sheet after the
recrystallization annealing process and the galvanized steel sheet
after the plating process or after the alloying process may be
subjected to a temper rolling at a rolling reduction of not more
than 10% for correcting the shape and adjusting the surface
roughness and the like. Furthermore, the cold rolled steel sheet
according to the invention can be applied as not only a cold rolled
steel sheet for the working but also a blank of a surface treated
steel sheet for the working. As the surface treated steel sheet for
the working are mentioned tin-plated steel sheets, porcelain
enamels and so on in addition to the aforementioned galvanized
steel sheets (including alloyed sheets). There is no problem even
when they are subjected to a treatment such as resin or fat
coating, various paintings, electroplating or the like. Moreover,
the galvanized steel sheet according to the invention may be
subjected to a special treatment after the galvanization in order
to improve the chemical conversion property, weldability, press
formability, corrosion resistance and the like.
BRIEF DESCRIPTION OF THE DRAWINGS
[0121] FIG. 1a is a graph showing an influence of V and C contents
in steel upon a Lankford value (r-value).
[0122] FIG. 1b is a graph showing an influence of V and C contents
in steel upon a yield ratio (YR=yield stress(YS)/tensile
stress(TS).times.100(%)).
[0123] FIG. 2a is a graph showing an influence of a relationship
among Nb, Ti and V contents upon a tensile strength (TS) in the
high-strength dual-phase cold rolled steel sheets of TS: 590 MPa
grade and TS: 780 MPa grade.
[0124] FIG. 2b is a graph showing an influence of a relationship
among Nb, Ti and V contents upon a Lankford value (r-value) in the
high-strength dual-phase cold rolled steel sheets of TS: 590 MPa
grade and TS: 780 MPa grade.
BEST MODE FOR CARRYING OUT THE INVENTION
[0125] Each of molten steels having compositions shown in Tables
1-4 is made in a converter and subjected to a continuous casting
process to obtain a slab. In this case, each of the slabs having
the compositions shown in Tables 1 and 2 is prepared for the
purpose of experiments with respect to the cold rolled steel sheet,
and each of the slabs having the compositions shown in Tables 3 and
4 is prepared for the purpose of experiments with respect to the
galvanized steel sheet. Especially, the slabs shown in Tables 2 and
4 are prepared for the purpose of obtaining the cold rolled steel
sheet and galvanized steel sheet of TS: not less than 780 MPa,
respectively. Then, the steel slab is heated to 1150.degree. C. and
subjected to a hot rolling under conditions of a finish deformation
temperature: 900.degree. C. and a coiling temperature: 650.degree.
C. at a hot rolling step to obtain a hot rolled steel strip having
a thickness of 4.0 mm. Subsequently, the hot rolled steel strip is
pickled and subjected to a cold rolling at a rolling reduction of
70% at a cold rolling step to obtain a cold rolled steel strip or a
cold rolled sheet having a thickness of 1.2 mm. Next, each of the
cold rolled steel sheets in Tables 1 and 2 is subjected to a
recrystallization annealing at an annealing temperature shown in
Tables 5 and 6 in a continuous annealing line. The thus obtained
cold rolled sheet is further subjected to a temper rolling at a
rolling reduction of 0.8%. With respect to the galvanized steel
sheets, each of the cold rolled sheets in Tables 3 and 4 is
subjected to a recrystallization annealing at an annealing
temperature shown in Tables 7 and 8 and further to a galvanizing in
a galvanizing bath having an Al concentration of 0.13% in a
continuous galvanization line. Moreover, with respect to a part of
steel sheets (Steel sheet Nos. 52, 68, 69 and 70 in Table 7), the
steel sheet after the cold rolling is subjected to an annealing at
830.degree. C. in a continuous annealing line and then pickled and
annealed and galvanized at a galvanizing bath temperature of
480.degree. C. under an Al concentration in the bath of 0.13% in a
continuous galvanization line and further the thus obtained steel
strip (galvanized steel sheet) is subjected to a temper rolling at
a rolling reduction of 0.8%. With respect to the steel sheets 75
and 77 in Table 7, they are subjected to an alloying treatment at
an alloying temperature of 520.degree. C. after the
galvanization.
[0126] A test piece is cut out from the obtained steel strip and a
microstructure thereof with respect to a section (C section)
perpendicular to the rolling direction is imaged by using an
optical microscope or a scanning electron microscope to measure a
structure ratio of ferrite phase as a primary phase and a kind and
a structure ratio of a secondary phase by using an image analysis
device. In this case, a specimen for observing the microstructure
is subjected to a mirror-like polishing and an etching with an
alcohol solution containing 2% HNO.sub.3 and then used for the
observation. And also, a tensile test piece of JIS No. 5 is cut out
from the steel strip and subjected to a tensile test according to
the definition of JIS Z 2241 to measure a yield stress (YS), a
tensile strength (TS), an elongation (E1), a yield ratio (YR) and a
Lankford value (r-value). These results are shown in Tables
5-8.
1TABLE 1(a) Steel Chemical composition (mass %) No. C Si Mn P S Al
N V Nb Ti Cr Mo Cu Ni 1-A 0.030 0.02 1.55 0.01 0.004 0.032 0.002
0.132 -- -- -- -- -- -- 1-B 0.028 0.02 1.48 0.01 0.001 0.032 0.002
0.105 0.042 -- -- 0.15 -- -- 1-C 0.032 0.03 1.72 0.01 0.005 0.028
0.002 0.085 0.035 0.035 0.05 -- -- -- 1-D 0.020 0.02 1.63 0.01
0.005 0.033 0.002 0.065 -- -- -- -- 0.12 0.08 1-E 0.031 0.02 1.56
0.01 0.006 0.033 0.002 0.122 0.045 -- -- 0.18 -- -- 1-F 0.029 0.02
1.48 0.01 0.003 0.032 0.002 0.210 0.115 0.125 -- -- -- -- 1-G 0.032
0.02 1.65 0.01 0.004 0.032 0.002 0.045 -- -- -- -- -- -- 1-H 0.020
0.22 2.02 0.06 0.004 0.032 0.002 0.132 -- -- -- -- -- 1-I 0.022
0.52 1.85 0.03 0.001 0.032 0.002 0.105 0.042 -- -- 0.15 -- -- 1-J
0.028 0.33 1.72 0.01 0.005 0.028 0.002 0.085 0.035 0.035 0.05 -- --
-- 1-K 0.011 0.21 1.53 0.01 0.003 0.028 0.002 0.032 0.030 -- -- --
-- -- 1-L 0.022 0.52 1.52 0.01 0.002 0.033 0.002 0.125 -- 0.022 --
-- -- -- 1-M 0.019 0.53 1.43 0.05 0.001 0.032 0.002 0.105 -- --
0.05 0.15 -- -- Transformation point Steel (.degree. C.) No.
X*.sup.1 Y*.sup.2 Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 1-A 1.04 -- --
725 860 Acceptable example 1-B -- 1.27 0.44 705 855 Acceptable
example 1-C -- 1.45 1.33 710 850 Acceptable example 1-D 0.76 -- --
715 855 Acceptable example 1-E -- 1.30 0.40 705 855 Acceptable
example 1-F -- 4.88 1.26 725 855 Comparative example 1-G 0.33 -- --
715 850 Comparative example 1-H 1.55 -- -- 725 860 Acceptable
example 1-I -- 1.62 0.44 705 865 Acceptable example 1-J -- 1.66
1.33 710 860 Acceptable example 1-K -- 1.39 1.03 710 860 Acceptable
example 1-L -- 1.84 0.37 715 865 Acceptable example 1-M 1.30 -- --
710 850 Acceptable example (Note) *.sup.1X = (V/51)/(C/12) *.sup.2Y
= (V/51 + 2 .times. Nb/93 + 2 .times. Ti/48)/(C/12) *.sup.3Z = (2
.times. Nb/93 + 2 .times. Ti/48)/(V/51)
[0127]
2TABLE 1(b) Steel Chemical composition (mass %) No. C Si Mn P S Al
N V Nb Ti Cr Mo Cu Ni 1-N 0.021 0.33 1.72 0.06 0.003 0.030 0.002
0.115 -- -- 0.05 0.15 0.15 0.15 1-O 0.020 0.41 2.02 0.02 0.002
0.029 0.002 0.072 0.042 0.010 0.05 0.15 0.10 0.10 1-P 0.007 0.35
1.76 0.01 0.005 0.029 0.002 0.073 -- -- -- -- -- -- 1-Q 0.112 0.33
1.74 0.01 0.003 0.028 0.002 0.352 -- -- -- -- -- -- 1-R 0.021 0.52
1.52 0.01 0.002 0.033 0.002 0.008 -- -- -- -- -- -- 1-S 0.023 0.53
1.43 0.05 0.001 0.032 0.002 0.622 -- -- -- -- -- -- 1-T 0.021 0.33
1.72 0.06 0.003 0.030 0.002 0.049 0.0005 -- -- -- -- -- 1-U 0.025
0.41 1.75 0.04 0.002 0.029 0.002 0.041 0.325 -- -- -- -- -- 1-V
0.019 0.35 1.76 0.05 0.001 0.032 0.002 0.052 -- 0.0005 -- -- -- --
1-W 0.023 0.33 1.72 0.06 0.003 0.030 0.002 0.033 -- 0.306 -- -- --
-- 1-X 0.018 0.02 1.48 0.01 0.003 0.032 0.002 0.030 0.001 0.001 --
-- -- -- 1-Y 0.021 0.02 1.65 0.01 0.004 0.032 0.002 0.329 -- -- --
-- -- -- Transformation point Steel (.degree. C.) No. X*.sup.1
Y*.sup.2 Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 1-N 1.29 -- -- 705 855
Acceptable example 1-O -- 1.64 0.93 715 850 Acceptable example 1-P
2.45 -- -- 714 882 Comparative example 1-Q 0.74 -- -- 714 859
Comparative example 1-R 0.09 -- -- 722 879 Comparative example 1-S
6.36 -- -- 723 972 Comparative example 1-T -- 0.56 0.01 714 910
Acceptable example 1-U -- 3.74 8.69 716 891 Comparative example 1-V
-- 0.66 0.02 714 906 Acceptable example 1-W -- 6.99 19.7 714 1033
Comparative example 1-X -- 0.43 0.07 708 863 Comparative example
1-Y 3.69 -- -- 706 886 Comparative example (Note) *.sup.1X =
(V/51)/(C/12) *.sup.2Y = (V/51 + 2 .times. Nb/93 + 2 .times.
Ti/48)/(C/12) *.sup.3Z = (2 .times. Nb/93 + 2 .times.
Ti/48)/(V/51)
[0128]
3TABLE 2 Steel Chemical composition (mass %) No. C Si Mn P S Al N V
Nb Ti Cr Mo Cu Ni 2-A 0.039 0.50 2.85 0.01 0.005 0.031 0.002 0.151
-- -- -- -- -- -- 2-B 0.038 0.75 2.52 0.01 0.001 0.035 0.002 0.088
0.121 -- -- 0.29 -- -- 2-C 0.042 0.74 2.53 0.01 0.006 0.033 0.002
0.092 0.110 0.152 0.09 -- -- -- 2-D 0.041 0.70 2.55 0.01 0.008
0.032 0.002 0.087 -- 0.064 -- -- 0.08 0.1 2-E 0.048 0.72 2.52 0.01
0.005 0.034 0.002 0.153 0.202 0.005 -- 0.31 -- -- 2-F 0.040 0.77
2.55 0.01 0.007 0.036 0.002 0.524 0.193 0.262 -- -- -- -- 2-G 0.038
0.73 2.56 0.01 0.006 0.033 0.002 0.040 0.011 0.009 -- -- -- -- 2-H
0.043 0.95 2.95 0.05 0.005 0.032 0.002 0.095 0.002 0.119 -- 0.27 --
-- 2-I 0.042 0.82 2.78 0.04 0.009 0.035 0.002 0.141 0.045 0.053 --
-- -- -- 2-J 0.048 0.91 2.73 0.04 0.006 0.036 0.002 0.033 0.185
0.155 0.13 0.13 0.12 0.11 2-K 0.042 0.82 2.78 0.04 0.009 0.035
0.002 0.008 -- -- -- -- -- 2-L 0.038 0.91 2.73 0.04 0.006 0.036
0.002 0.522 -- -- -- -- -- 2-M 0.038 0.76 2.57 0.03 0.001 0.034
0.002 0.087 0.0005 -- -- -- -- -- 2-N 0.039 0.76 2.55 0.03 0.001
0.035 0.002 0.032 0.056 -- -- -- -- -- 2-O 0.042 0.73 2.49 0.03
0.001 0.036 0.002 0.092 0.382 -- -- -- -- -- 2-P 0.043 0.75 2.52
0.03 0.001 0.035 0.002 0.088 0.453 -- -- -- -- -- 2-Q 0.041 0.70
2.55 0.03 0.002 0.032 0.002 0.098 -- 0.0005 -- -- -- -- 2-R 0.038
0.71 2.58 0.03 0.002 0.030 0.002 0.025 -- 0.037 -- -- -- -- 2-S
0.039 0.74 2.57 0.04 0.002 0.030 0.002 0.079 -- 0.186 -- -- -- --
2-T 0.042 0.71 2.51 0.02 0.002 0.029 0.002 0.089 -- 0.356 -- -- --
-- Transformation point Steel (.degree. C.) No. X*.sup.1 Y*.sup.2
Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 2-A 0.91 -- -- 707 842
Acceptable example 2-B -- 1.37 1.51 718 868 Acceptable example 2-C
-- 3.00 4.82 719 915 Acceptable example 2-D -- 1.28 1.56 715 875
Acceptable example 2-E -- 1.89 1.52 717 870 Acceptable example 2-F
-- 7.60 1.47 718 972 Comparative example 2-G -- 0.44 0.78 717 851
Comparative example 2-H -- 1.92 2.68 717 931 Acceptable example 2-I
-- 1.70 1.15 717 900 Acceptable example 2-J -- 2.77 16.13 722 932
Acceptable example 2-K 0.04 -- -- 717 866 Comparative example 2-L
3.23 -- -- 720 923 Comparative example 2-M -- 0.54 0.01 718 868
Acceptable example 2-N -- 0.56 1.92 718 868 Acceptable example 2-O
-- 2.86 4.55 718 868 Acceptable example 2-P -- 3.20 5.65 718 868
Comparative example 2-Q -- 0.57 0.01 715 875 Acceptable example 2-R
-- 0.64 3.15 715 875 Acceptable example 2-S -- 2.86 5.00 715 875
Acceptable example 2-T -- 4.74 8.50 715 875 Comparative example
(Note) *.sup.1X = (V/51)/(C/12) *.sup.2Y = (V/51 + 2 .times. Nb/93
+ 2 .times. Ti/48)/(C/12) *.sup.3Z = (2 .times. Nb/93 + 2 .times.
Ti/48)/(V/51)
[0129]
4TABLE 3(a) Steel Chemical composition (mass %) No. C Si Mn P S Al
N V Nb Ti Cr Mo Cu Ni 3-A 0.028 0.02 1.55 0.01 0.003 0.034 0.002
0.121 -- -- -- -- -- -- 3-B 0.030 0.02 1.46 0.01 0.002 0.035 0.002
0.108 0.041 -- -- 0.16 -- -- 3-C 0.031 0.03 1.70 0.01 0.005 0.028
0.002 0.086 0.036 0.033 0.06 -- -- -- 3-D 0.021 0.02 1.65 0.01
0.005 0.034 0.002 0.068 -- -- -- -- 0.14 0.07 3-E 0.032 0.02 1.52
0.01 0.004 0.033 0.002 0.124 0.044 -- -- 0.15 -- -- 3-F 0.026 0.02
1.52 0.01 0.003 0.035 0.002 0.122 0.112 0.122 -- -- -- -- 3-G 0.032
0.02 1.62 0.01 0.005 0.032 0.002 0.042 -- -- -- -- -- -- 3-H 0.021
0.21 2.02 0.06 0.003 0.030 0.002 0.130 -- -- -- -- -- -- 3-I 0.024
0.52 1.88 0.04 0.001 0.032 0.002 0.105 0.033 -- 0.16 -- -- 3-J
0.026 0.32 1.72 0.01 0.004 0.026 0.002 0.088 0.035 0.032 0.08 -- --
-- 3-K 0.020 0.70 1.55 0.01 0.003 0.028 0.002 0.073 0.045 -- -- --
-- -- 3-L 0.012 0.21 1.51 0.01 0.002 0.033 0.002 0.055 -- 0.018 --
-- -- -- 3-M 0.018 0.50 1.56 0.03 0.004 0.035 0.002 0.108 -- --
0.05 0.15 -- -- Transformation point Steel (.degree. C.) No.
X*.sup.1 Y*.sup.2 Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 3-A 1.02 -- --
725 860 Acceptable example 3-B -- 1.20 0.42 705 855 Acceptable
example 3-C -- 1.48 1.27 710 850 Acceptable example 3-D 0.76 -- --
715 855 Acceptable example 3-E -- 1.27 0.39 705 855 Acceptable
example 3-F -- 4.56 3.13 725 855 Comparative example 3-G 0.31 -- --
715 850 Comparative example 3-H 1.46 -- -- 725 860 Acceptable
example 3-I -- 1.38 0.34 705 860 Acceptable example 3-J -- 1.76
1.21 710 860 Acceptable example 3-K -- 1.44 0.68 715 870 Acceptable
example 3-L -- 1.83 0.70 710 865 Acceptable example 3-M 1.41 -- --
710 860 Acceptable example (Note) *.sup.1X = (V/51)/(C/12) *.sup.2Y
= (V/51 + 2 .times. Nb/93 + 2 .times. Ti/48)/(C/12) *.sup.3Z = (2
.times. Nb/93 + 2 .times. Ti/48)/(V/51)
[0130]
5TABLE 3(b) Steel Chemical composition (mass %) No. C Si Mn P S Al
N V Nb Ti Cr Mo Cu Ni 3-N 0.020 0.39 1.73 0.05 0.001 0.031 0.002
0.110 -- -- 0.05 0.15 0.15 0.15 3-O 0.021 0.28 1.95 0.02 0.005
0.029 0.002 0.075 0.038 0.01 0.05 0.15 0.10 0.10 3-P 0.008 0.32
1.75 0.01 0.005 0.032 0.002 0.075 -- -- -- -- -- -- 3-Q 0.095 0.34
1.73 0.01 0.003 0.029 0.002 0.361 -- -- -- -- -- -- 3-R 0.023 0.49
1.54 0.01 0.002 0.030 0.002 0.007 -- -- -- -- -- -- 3-S 0.024 0.51
1.47 0.03 0.001 0.031 0.002 0.597 -- -- -- -- -- -- 3-T 0.022 0.35
1.75 0.05 0.003 0.029 0.002 0.109 0.0005 -- -- -- -- -- 3-U 0.023
0.44 1.78 0.04 0.003 0.027 0.002 0.065 0.319 -- -- -- -- -- 3-V
0.021 0.35 1.73 0.05 0.001 0.034 0.002 0.099 -- 0.0005 -- -- -- --
3-W 0.025 0.36 1.77 0.05 0.002 0.032 0.002 0.132 -- 0.321 -- -- --
-- 3-X 0.020 0.02 1.51 0.01 0.003 0.033 0.002 0.035 0.001 0.001 --
-- -- -- 3-Y 0.023 0.02 1.66 0.01 0.003 0.035 0.002 0.308 -- -- --
-- -- -- Transformation point Steel (.degree. C.) No. X*.sup.1
Y*.sup.2 Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 3-N 1.29 -- -- 705 865
Acceptable example 3-O -- 1.55 0.84 715 865 Acceptable example 3-P
2.21 -- -- 714 881 Comparative example 3-Q 0.89 -- -- 714 872
Comparative example 3-R 0.07 -- -- 722 874 Comparative example 3-S
5.85 -- -- 721 949 Comparative example 3-T -- 1.17 0.01 714 906
Acceptable example 3-U -- 4.24 5.38 717 891 Comparative example 3-V
-- 1.12 0.01 715 913 Acceptable example 3-W -- 7.66 5.17 715 1031
Comparative example 3-X -- 0.45 0.09 707 866 Comparative example
3-Y 3.15 -- -- 707 882 Comparative example (Note) *.sup.1X =
(V/51)/(C/12) *.sup.2Y = (V/51 + 2 .times. Nb/93 + 2 .times.
Ti/48)/(C/12) *.sup.3Z = (2 .times. Nb/93 + 2 .times.
Ti/48)/(V/51)
[0131]
6TABLE 4 Steel Chemical composition (mass %) No. C Si Mn P S Al N V
Nb Ti Cr Mo Cu Ni 4-A 0.038 0.48 2.88 0.01 0.004 0.033 0.002 0.158
-- -- -- -- -- -- 4-B 0.041 0.77 2.51 0.01 0.001 0.035 0.002 0.056
0.171 -- -- 0.31 -- -- 4-C 0.040 0.76 2.49 0.01 0.007 0.034 0.002
0.068 0.120 0.125 0.09 -- -- -- 4-D 0.038 0.72 2.54 0.01 0.009
0.033 0.002 0.085 -- 0.058 -- -- 0.08 0.07 4-E 0.049 0.74 2.53 0.01
0.006 0.036 0.002 0.039 0.075 0.005 -- 0.31 -- -- 4-F 0.039 0.75
2.55 0.01 0.007 0.035 0.002 0.183 0.191 0.260 -- -- -- -- 4-G 0.046
0.73 2.57 0.01 0.007 0.038 0.002 0.011 0.013 0.015 -- -- -- -- 4-H
0.039 0.93 2.95 0.05 0.004 0.039 0.002 0.016 0.003 0.108 -- 0.27 --
-- 4-I 0.041 0.80 2.80 0.05 0.009 0.033 0.002 0.138 0.042 0.065 --
-- -- -- 4-J 0.047 0.92 2.78 0.04 0.006 0.034 0.002 0.025 0.175
0.143 0.15 -- -- -- 4-K 0.043 0.84 2.76 0.04 0.007 0.034 0.002
0.007 -- -- -- -- -- 4-L 0.038 0.93 2.75 0.04 0.006 0.035 0.002
0.553 -- -- -- -- -- 4-M 0.042 0.80 2.65 0.02 0.003 0.031 0.002
0.096 0.0005 -- -- -- -- -- 4-N 0.041 0.81 2.68 0.02 0.003 0.030
0.002 0.029 0.065 -- -- -- -- -- 4-O 0.043 0.78 2.67 0.03 0.003
0.029 0.002 0.087 0.295 -- -- -- -- -- 4-P 0.041 0.77 2.69 0.02
0.002 0.030 0.002 0.079 0.521 -- -- -- -- -- 4-Q 0.039 0.76 2.74
0.03 0.003 0.031 0.002 0.105 -- 0.0005 -- -- -- -- 4-R 0.043 0.79
2.74 0.02 0.004 0.033 0.002 0.035 -- 0.042 -- -- -- -- 4-S 0.040
0.80 2.75 0.03 0.002 0.032 0.002 0.087 -- 0.182 -- -- -- -- 4-T
0.038 0.81 2.77 0.02 0.003 0.032 0.002 0.089 -- 0.290 -- -- -- --
Transformation point Steel (.degree. C.) No. X*.sup.1 Y*.sup.2
Z*.sup.3 A.sub.c1 A.sub.c3 Remarks 4-A 0.98 -- -- 706 842
Acceptable example 4-B -- 1.40 3.35 719 865 Acceptable example 4-C
-- 2.74 5.84 720 905 Acceptable example 4-D -- 1.29 1.45 716 876
Acceptable example 4-E -- 0.63 2.38 717 859 Acceptable example 4-F
-- 5.70 4.16 718 971 Comparative example 4-G -- 0.29 4.19 717 851
Comparative example 4-H -- 1.50 14.55 717 922 Acceptable example
4-I -- 1.85 1.33 716 909 Acceptable example 4-J -- 2.61 19.83 723
923 Acceptable example 4-K 0.04 -- -- 718 867 Comparative example
4-L 3.42 -- -- 721 924 Comparative example 4-M -- 0.54 0.01 719 865
Acceptable example 4-N -- 0.58 2.46 719 865 Acceptable example 4-O
-- 2.25 3.72 719 865 Acceptable example 4-P -- 3.73 7.23 719 865
Comparative example 4-Q -- 0.64 0.01 716 876 Acceptable example 4-R
-- 0.68 2.55 716 876 Acceptable example 4-S -- 2.79 4.45 716 876
Acceptable example 4-T -- 4.37 6.92 716 876 Comparative example
(Note) *.sup.1X = (V/51)/(C/12) *.sup.2Y = (V/51 + 2 .times. Nb/93
+ 2 .times. Ti/48)/(C/12) *.sup.3Z = (2 .times. Nb/93 + 2 .times.
Ti/48)/(V/51)
[0132]
7TABLE 5(a) Cold rolling Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase cold rolled steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 1 1-A 830 92 M 8 8 330 600 31
55 1.8 Invention example 2 1-B 830 90 M 10 10 330 610 30 54 1.8
Invention example 3 1-B 980 0 P, B, M 15 100 650 720 22 90 0.9
Comparative example 4 1-B 680 100 -- 0 0 450 530 29 85 0.8
Comparative Example 5 1-C 830 92 M 8 8 340 600 31 57 1.8 Invention
example 6 1-D 830 90 M 10 10 330 610 30 54 1.4 Invention example 7
1-E 830 92 M 8 8 310 570 33 54 1.7 Invention example 8 1-F 830 100
-- 0 0 510 600 27 85 1.8 Comparative example 9 1-G 830 93 M 7 7 330
610 31 54 0.8 Comparative example 10 1-H 850 92 M 8 8 350 630 29 56
1.9 Invention example 11 1-I 850 93 M 7 7 330 620 30 53 1.9
Invention example 12 1-J 850 92 M 8 8 330 610 33 54 1.8 Invention
example 13 1-K 830 92 M 8 8 245 450 38 54 1.9 Invention example 14
1-L 830 93 M 7 7 330 605 30 55 1.8 Invention example (Note)
*.sup.1F is abbreviation of ferrite phase, M is abbreviation of
matensite phase, P is abbreviation of perlite phase and B is
abbreviation of beinite phase.
[0133]
8TABLE 5(b) Cold rolling Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase cold rolled steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 15 1-M 830 92 M 8 8 340 620 30
55 1.7 Invention example 16 1-N 830 93 M 7 7 320 600 31 53 1.7
Invention example 17 1-O 830 92 M, B 6 8 340 625 29 54 1.8
Invention example 18 1-P 830 100 -- 0 0 425 520 34 82 1.9
Comparative Example 19 1-Q 830 65 M 35 35 395 670 29 59 0.8
Comparative example 20 1-R 850 69 M 31 31 370 620 30 60 0.8
Comparative example 21 1-S 850 100 -- 0 0 495 615 30 80 1.7
Comparative example 22 1-T 850 92 M 8 8 355 575 32 62 1.7 Invention
example 23 1-U 850 100 -- 0 0 470 580 31 81 1.8 Comparative example
24 1-V 830 91 M 9 9 350 570 32 61 1.7 Invention example 25 1-W 850
100 -- 0 0 480 595 31 81 1.8 Comparative example 26 1-X 830 72 M 28
28 350 560 31 63 0.8 Comparative example 27 1-Y 830 100 -- 0 0 475
590 30 81 1.7 Comparative example (Note) *.sup.1F is abbreviation
of ferrite phase, M is abbreviation of matensite phase, P is
abbreviation of perlite phase and B is abbreviation of beinite
phase.
[0134]
9TABLE 6(a) Cold rolling Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase cold rolled steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 28 2-A 780 90 M 10 10 560 825
19 68 1.1 Invention example 29 2-B 780 87 M 13 13 550 810 19 68 1.3
Invention example 30 2-B 950 0 P, B, M 19 100 740 860 16 86 0.7
Comparative example 31 2-B 680 100 -- 0 0 625 770 22 81 0.8
Comparative Example 32 2-C 750 88 M 12 12 540 805 20 67 1.3
Invention example 33 2-D 760 88 M 12 12 545 810 19 67 1.2 Invention
example 34 2-E 770 87 M 13 13 550 820 20 67 1.3 Invention example
35 2-F 780 100 -- 0 0 660 830 19 80 1.4 Comparative example 36 2-G
780 69 M 31 31 540 820 20 66 0.7 Comparative example 37 2-H 760 81
M 19 19 620 930 15 67 1.3 Invention example 38 2-I 780 83 M 17 17
590 860 17 69 1.1 Invention example (Note) *.sup.1F is abbreviation
of ferrite phase, M is abbreviation of matensite phase, P is
abbreviation of perlite phase and B is abbreviation of beinite
phase.
[0135]
10TABLE 6(b) Cold rolling Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase cold rolled steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 39 2-J 780 87 M 13 13 445 660
27 67 1.4 Invention example 40 2-K 760 68 M 32 32 570 850 18 67 0.8
Comparative example 41 2-L 780 100 -- 0 0 690 835 19 83 1.3
Comparative example 42 2-M 780 85 M 15 15 525 805 20 65 1.1
Invention example 43 2-N 760 88 M 12 12 530 800 20 66 1.3 Invention
example 44 2-O 780 90 M 10 10 525 790 21 66 1.3 Invention example
45 2-P 780 100 -- 0 0 650 795 21 82 1.3 Comparative example 46 2-Q
760 87 M 13 13 540 810 19 67 1.1 Invention example 47 2-R 760 88 M
12 12 545 815 15 67 1.3 Invention example 48 2-S 780 90 M 10 10 540
810 19 67 1.3 Invention example 49 2-T 780 100 -- 0 0 665 785 20 85
1.4 Comparative example (Note) *.sup.1F is abbreviation of ferrite
phase, M is abbreviation of matensite phase, P is abbreviation of
perlite phase and B is abbreviation of beinite phase.
[0136]
11TABLE 7(a) Galvanizing Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase galvanized steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 50 3-A 830 92 M 8 8 330 610 31
54 1.7 Invention example 51 3-B 830 90 M 10 10 330 620 30 53 1.7
Invention example 52 3-B 830 92 M 8 8 350 630 30 56 1.6 Invention
example 53 3-B 980 0 P, B, M 12 100 660 720 22 92 0.9 Comparative
Example 54 3-B 680 100 -- 0 0 460 540 28 85 0.8 Comparative example
55 3-C 830 90 M 10 10 340 610 31 56 1.7 Invention example 56 3-D
830 92 M 8 8 340 620 30 55 1.4 Invention example 57 3-E 830 94 M 6
6 320 580 32 55 1.6 Invention example 58 3-F 830 100 -- 0 0 510 600
27 85 1.7 Comparative example 59 3-G 830 92 M 8 8 330 610 30 54 0.8
Comparative example 60 3-H 850 93 M 7 7 340 630 30 54 1.8 Invention
example 61 3-I 850 92 M 8 8 340 620 31 55 1.8 Invention example 62
3-J 850 92 M 8 8 320 610 31 52 1.7 Invention example 63 3-K 830 92
M, B 6 8 330 610 30 54 1.6 Invention example 64 3-L 830 92 M 8 8
248 450 37 55 1.7 Invention example 65 3-M 830 93 M 7 7 340 620 30
55 1.6 Invention example (Note) *.sup.1F is abbreviation of ferrite
phase, M is abbreviation of matensite phase, P is abbreviation of
perlite phase and B is abbreviation of beinite phase.
[0137]
12TABLE 7(b) Galvanizing Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase galvanized steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 66 3-N 830 92 M 8 8 320 600 31
53 1.6 Invention example 67 3-O 830 93 M 7 7 340 625 29 54 1.7
Invention example 68 3-H 830 92 M 8 8 340 620 30 55 1.8 Invention
example 69 3-K 830 93 M 7 7 320 600 31 53 1.6 Invention example 70
3-M 830 92 M 8 8 320 610 31 52 1.6 Invention example 71 3-P 830 100
-- 0 0 420 510 34 82 1.8 Comparative example 72 3-Q 830 66 M 34 34
390 670 27 58 0.8 Comparative example 73 3-R 850 68 M 32 32 385 615
30 63 0.8 Comparative example 74 3-S 850 100 -- 0 0 500 605 31 83
1.6 Comparative example 75 3-T 850 91 M 9 9 350 580 31 60 1.7
Invention example 76 3-U 850 100 -- 0 0 480 575 32 83 1.6
Comparative example 77 3-V 830 91 M 9 9 340 580 31 59 1.7 Invention
example 78 3-W 850 100 -- 0 0 490 600 30 82 1.7 Comparative example
79 3-X 830 70 M 30 30 340 565 32 60 0.8 Comparative example 80 3-Y
830 100 -- 0 0 490 600 30 82 1.7 Comparative example (Note)
*.sup.1F is abbreviation of ferrite phase, M is abbreviation of
matensite phase, P is abbreviation of perlite phase and B is
abbreviation of beinite phase.
[0138]
13TABLE 8(a) Galvanizing Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase galvanized steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 81 4-A 780 91 M 9 9 560 815 19
69 1.1 Invention example 82 4-B 780 89 M 11 11 555 805 19 69 1.4
Invention example 83 4-B 950 0 P,B,M 21 100 735 850 16 86 0.8
Comparative example 84 4-B 680 100 -- 0 0 620 760 22 82 0.8
Comparative Example 85 4-C 4 89 M 11 11 545 800 20 68 1.3 Invention
example 86 4-D 760 88 M 12 12 550 805 19 68 1.4 Invention example
87 4-E 770 90 M 10 10 550 810 20 68 1.3 Invention example 88 4-F
780 100 -- 0 0 675 815 19 83 1.5 Comparative example 89 4-G 780 92
M 8 8 550 810 20 68 0.8 Comparative example 90 4-H 760 83 M 17 17
635 935 15 68 1.3 Invention example 91 4-I 780 85 M 15 15 590 855
17 69 1.1 Invention example (Note) *.sup.1F is abbreviation of
ferrite phase, M is abbreviation of matensite phase, P is
abbreviation of perlite phase and B is abbreviation of beinite
phase.
[0139]
14TABLE 8(b) Galvanizing Annealing Microstructure Mechanical
properties of temperature in Ferrite Second phase galvanized steel
sheet Steel continuous phase Area ratio of Area ratio Tensile
properties sheet Steel annealing line Area martensite of second YS
TS El YR No. No. (.degree. C.) ratio (%) Kind*.sup.1 (%) phase (%)
(MPa) (MPa) (%) (%) r-value Remarks 92 4-J 780 85 M 15 15 440 665
25 68 1.4 Invention example 93 4-K 760 67 M 33 33 560 860 18 65 0.8
Comparative example 94 4-L 780 100 -- 0 0 695 840 19 83 1.4
Comparative example 95 4-M 780 86 M 14 14 510 810 20 63 1.1
Invention example 96 4-N 760 89 M 11 11 525 800 20 66 1.3 Invention
example 97 4-O 780 89 M 11 11 525 795 20 66 1.3 Invention example
98 4-P 780 100 -- 0 0 660 805 20 82 1.4 Comparative example 99 4-Q
760 87 M 13 13 525 810 19 65 1.1 Invention example 100 4-R 760 86 M
14 14 530 810 19 65 1.2 Invention example 101 4-S 780 89 M 11 11
540 820 18 66 1.3 Invention example 102 4-T 780 100 -- 0 0 660 790
20 84 1.3 Comparative example (Note) *.sup.1F is abbreviation of
ferrite phase, M is abbreviation of matensite phase, P is
abbreviation of perlite phase and B is abbreviation of beinite
phase.
[0140] As seen from the results shown in Tables 5 and 6, the cold
rolled steel sheets in all invention examples have a low yield
stress (YS), a high elongation (E1) and a low yield ratio (YR) and
further indicate a high r-value and are excellent in the deep
drawability, and have a tensile strength (TS) of not less than 440
MPa. On the contrary, in the comparative examples being outside the
range of the invention, the yield stress (YS) is high, the
elongation (E1) is low, or the r-value is low. Particularly, the
somewhat lowering of the r-value accompanied with the
high-strengthening is observed in the high-strength steel sheets of
TS: not less than 780 MPa shown in Table 6, for example, the steel
sheet No. 28 produced by using the steel No. 2-A containing V and
no Nb and Ti and the steel sheet No. 38 produced by using the steel
No. 2-I containing V, Nb and Ti and satisfying a relationship of
0.5.times.C/12.ltoreq.(V/51+2.-
times.Nb/93+2.times.Ti/48).ltoreq.3.times.C/12 but satisfying a
relationship of (2.times.Nb/93+2.times.Ti/48)/(V/51)<0.5. On the
other hand, the r-value is improved in the steel sheet Nos. 29, 32,
33 and 34 produced by using the steel Nos. 2-B, 2-C, 2-D and 2-E
containing V, Nb and Ti and satisfying both relationships of
0.5.times.C/12.ltoreq.(V/51+2-
.times.Nb/93+2.times.Ti/48).ltoreq.3.times.C/12 and
1.5.ltoreq.(2.times.Nb/93+2.times.Ti/48)/(V/51).ltoreq.15.
[0141] And also, the results obtained with respect to the
galvanized steel sheets are shown in Tables 7 and 8. Even in these
galvanized steel sheets, the results similar to those of the above
cold rolled steel sheets are obtained.
[0142] In the steel sheet according to the invention, excellent
properties are obtained even by the production process conducting
the galvanization.
[0143] Industrial Applicability
[0144] The invention develops an industrially remarkable effect
that the high-strength cold rolled steel sheet and galvanized steel
sheet having an excellent deep drawability can be produced stably.
When the cold rolled steel sheet and the galvanized steel sheet
according to the invention are applied to vehicle parts, there are
effects that the press forming is easy and they can sufficiently
contribute to reduce the weight of the vehicle body.
* * * * *