U.S. patent application number 10/031426 was filed with the patent office on 2003-01-23 for cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same.
Invention is credited to Kami, Chikara, Tosaka, Akio, Yamazaki, Takuya.
Application Number | 20030015263 10/031426 |
Document ID | / |
Family ID | 27343520 |
Filed Date | 2003-01-23 |
United States Patent
Application |
20030015263 |
Kind Code |
A1 |
Kami, Chikara ; et
al. |
January 23, 2003 |
Cold rolled steel sheet and galvanized steel sheet having strain
aging hardening property and method for producing the same
Abstract
An object of the present invention is to provide a cold-rolled
steel sheet and an alloyed hot-dip galvanized steel sheet in which
tensile strength is effectively increased by press forming and heat
treatment while maintaining excellent deep drawability in press
forming. Specifically, a steel composition contains less than 0.01%
of C, 0.005 to 1.0% of Si, 0.01 to 1.0% of Mn, 0.005 to 0.050% of
Nb, 0.005 to 0.030% of Al, 0.005 to 0.040% of N, 0.0005 to 0.0015%
of B, 0.05% or less of P, and 0.01% or less of S, the balance
substantially composed of Fe, in which the following equations (1)
and (2) are satisfied: N %.gtoreq.0.0015+14/93.multidot.Nb
%+14/27.multidot.Al %+14/11.multidot.B % (1) C
%.ltoreq.(12/93).multidot.Nb % (2)
Inventors: |
Kami, Chikara; (Chiba,
JP) ; Tosaka, Akio; (Chiba, JP) ; Yamazaki,
Takuya; (Chiba, JP) |
Correspondence
Address: |
SCHNADER HARRISON SEGAL & LEWIS, LLP
1600 MARKET STREET
SUITE 3600
PHILADELPHIA
PA
19103
|
Family ID: |
27343520 |
Appl. No.: |
10/031426 |
Filed: |
January 18, 2002 |
PCT Filed: |
February 14, 2001 |
PCT NO: |
PCT/JP01/01004 |
Current U.S.
Class: |
148/603 ;
148/320 |
Current CPC
Class: |
C21D 8/0273 20130101;
C21D 9/46 20130101; C22C 38/001 20130101; Y10T 428/12799 20150115;
C22C 38/04 20130101; C22C 38/004 20130101; C21D 8/0436 20130101;
C22C 38/06 20130101; C21D 2211/008 20130101; C22C 38/12 20130101;
C21D 9/48 20130101; C22C 38/02 20130101; C21D 8/0473 20130101; Y10T
428/12493 20150115; C21D 2211/005 20130101; C21D 1/185 20130101;
C21D 8/0226 20130101; C21D 8/0236 20130101; C21D 8/0426
20130101 |
Class at
Publication: |
148/603 ;
148/320 |
International
Class: |
C21D 008/02 |
Foreign Application Data
Date |
Code |
Application Number |
May 26, 2000 |
JP |
2000-156274 |
Jun 28, 2000 |
JP |
2000-193717 |
Feb 14, 2001 |
JP |
2000-328924 |
Claims
What is claimed is:
1. A cold-rolled steel sheet having excellent strain age
hardenability comprising a composition, by mass %, comprising: C:
0.15% or less; Si: 1.0% or less; Mn: 2.0% or less; P: 0.1% or less;
S: 0.01% or less; Al: 0.005 to 0.030%; and N: 0.0050 to 0.0400%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010%
or more, and the balance is composed of Fe and inevitable
impurities.
2. A cold-rolled steel sheet having excellent strain age
hardenability according to claim 1, comprising a composition, by
mass %, further comprising at least one of the following groups a
to d: Group a: at least one of Cu, Ni, Cr and Mo in a total of 1.0%
or less; Group b: at least one of Nb, Ti and V in a total of 0.1%
or less; Group c: 0.0030% or less of B; and Group d: one or both of
Ca and REM in a total of 0.0010 to 0.010%.
3. A cold-rolled steel sheet having excellent strain age
hardenability comprising a composition, by mass %, comprising: C:
less than 0.01%; Si: 0.005 to 1.0%; Mn: 0.01 to 1.5%; P: 0.1% or
less; S: 0.01% or less; Al: 0.005 to 0.030%; and N: 0.005 to
0.040%; wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, and the balance is composed of Fe and inevitable
impurities.
4. A cold-rolled steel sheet having excellent strain age
hardenability according to claim 3, wherein the composition, by
mass %, preferably further comprises: B: 0.0001 to 0.0030%; and Nb:
0.005 to 0.050%; wherein the ranges of B and Nb satisfy the
following equations (1) and (2), and the balance is substantially
composed of Fe: N %.gtoreq.0.0015+14/93.mult- idot.Nb
%+14/27.multidot.Al %+14/11.multidot.B % (1)
C.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
5. A cold-rolled steel sheet having excellent strain age
hardenability according to claim 3 or 4, wherein the composition,
by mass %, further comprises at least one of Cu, Ni and Mo in a
total of 1.0% or less.
6. A cold-rolled steel sheet having excellent strain age
hardenability according to claims 1 to 5, wherein the steel sheet
has a crystal grain diameter of 20 .mu.m or less.
7. A cold-rolled steel sheet having excellent strain age
hardenability according to claims 1 to 6, wherein strength after
forming is increased by 60 MPa or more by heat treatment in a low
temperature region of 120 to 200.degree. C.
8. An electro-galvanized, hot-dip galvanized or alloyed hot-dip
galvanized steel sheet having excellent strain age hardenability
comprising a coated layer formed on the surface of a cold-rolled
steel sheet according to claims 1 to 7 by electro-galvanization,
hot-dip galvanization, or alloying hot-dip galvanization.
9. A method of producing a cold-rolled steel sheet having excellent
strain age hardenability comprising hot-rolling a steel slab under
conditions in which the steel slab is cooled immediately after the
end of finish rolling and coiled at a coiling temperature of 400 to
800.degree. C., cold-rolling the hot-rolled sheet with a rolling
reduction ratio of 60 to 95%, and then recrystallizing and
annealing the cold-rolled sheet at a temperature of 650 to
900.degree. C., wherein the steel slab has a composition, by mass
%, comprising: C: less than 0.01%; Si: 0.005 to 1.0%; Mn: 0.01 to
1.5%; P: 0.1% or less; S: 0.01% or less; Al: 0.005 to 0.030%; and
N: 0.005 to 0.040%; wherein N/Al is 0.30 or more, and the balance
is substantially composed of Fe.
10. A method of producing a cold-rolled steel sheet having
excellent strain age hardenability according to claim 9, wherein
the composition, by mass %, further comprises: B: 0.0001 to
0.0030%; and Nb: 0.005 to 0.050%; wherein the ranges of B and Nb
satisfy the following equations (1) and (2), and the balance is
substantially composed of Fe: N %.gtoreq.0.0015+14/93.multidot.Nb
%+14/27.multidot.Al %+14/11.multidot.B % (1) C
%.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
11. A method of producing a cold-rolled steel sheet having
excellent strain age hardenability according to claim 9 or 10,
wherein in the heating-up step in the recrystallization annealing,
the temperature is increased at a rate of 1 to 20.degree. C./s in
the temperature region from 500.degree. C. to the recrystallization
temperature.
12. A method of producing an galvannealed steel sheet having
excellent strain age hardenability comprising hot-dip galvanization
and then heat alloying after recrystallization and annealing
according to claims 9 to 11.
13. A cold-rolled deep drawing steel sheet having excellent strain
age hardenability comprising a composition, by mass %, comprising:
C: 0.01% or less; Si: 1.0% or less; Mn: 0.01 to 1.5%; P: 0.1% or
less; S: 0.01% or less; Al: 0.005 to 0.020%; and N: 0.0050 to
0.040%; wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, the balance is composed of Fe and inevitable
impurities, and Ts x r value is 750 MPa or more.
14. A cold-rolled deep drawing steel sheet having excellent strain
age hardenability according to claim 13, wherein the composition,
by mass %, further comprises: B: 0.0001 to 0.0030%; and Nb: 0.005
to 0.050%; wherein the balance is composed of Fe and inevitable
impurities, TS.times.r value is 750 MPa or more, and the ranges of
B and Nb satisfy the following equations (1) and (2),: N
%.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1) C %.ltoreq.0.5.multidot.(12-
/93).multidot.Nb % (2)
15. A cold-rolled deep drawing steel sheet having excellent strain
age hardenability according to claim 13, wherein the composition,
by mass %, further comprises at least one of the following: B:
0.0001 to 0.0030%; Nb: 0.005 to 0.050%; Ti: 0.005 to 0.070%; and V:
0.005 to 0.10%; wherein N/(Al+Nb+Ti+V+B) is 0.30 or more, the
amount of dissolved N is 0.0010% or more, the balance is composed
of Fe and inevitable impurities, and TS x r value is 750 MPa or
more.
16. A method of producing a cold-rolled deep drawing steel sheet
having excellent strain age hardenability comprising heating a
steel raw material to 950.degree. C. or more, roughly rolling the
raw material so that the finisher delivery temperature is Ar.sub.3
to 1000.degree. C., finish-rolling the material while lubricating
it in the temperature region of 600.degree. C. to Ar.sub.3, coiling
the rolled sheet in which the total reduction by rolling starting
from rough rolling to finish rolling is 80% or more,
recrystallizing and annealing the hot-rolled sheet, cold-rolling
the rolled sheet with a rolling reduction of 60 to 95%, and then
recrystallizing and annealing the resultant cold-rolled sheet,
wherein the steel raw material has a composition, by mass %,
comprising: C: less than 0.01%; Si: 0.005 to 1.0%; Mn: 0.01 to
1.5%; P: 0.1% or less; S: 0.01% or less; Al: 0.005 to 0.030%; N:
0.005 to 0.040%; and at least one of the following: B: 0.0003 to
0.0030%; Nb: 0.005 to 0.050%; Ti: 0.005 to 0.070%; and V: 0.005 to
0.10%; wherein N/(Al+Nb+Ti+V+B) is 0.30 or more.
17. A high-tensile-strength cold-rolled steel sheet having
excellent moldability, strain age hardenability and natural aging
resistance comprising a composition, by mass %, comprising: C:
0.0015 to 0.025%; Si: 1.0% or less; Mn: 2.0% or less; P: 0.1% or
less; S: 0.02% or less; Al: 0.02% or less; N: 0.0050 to 0.0250%;
and one or both of the following: B: 0.0001 to 0.0050%; and Nb:
0.002 to 0.050%; wherein N/Al is 0.30 or more, the amount of
dissolved N is 0.0010% or more, the balance is composed of Fe and
inevitable impurities, the structure is composed of an acicular
ferrite phase at an area ratio of 5% or more and a ferrite phase
having an average crystal grain diameter of 20 .mu.m or less, and
the r value is 1.2 or more.
18. A cold-rolled steel sheet according to claim 17, wherein the
composition further comprises at least one of the following groups
a to c: Group a: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less; Group b: one or both of Ti and V in a total of 0.1%
or less; and Group c: one or both of Ca and REM in a total of
0.0010 to 0.010%.
19. A method of producing a high-tensile-strength cold-rolled steel
sheet having a r value of 1.2 or more, and excellent moldability,
strain age hardenability and natural aging resistance comprising:
the hot-rolling step of roughly rolling a steel slab by heating to
a slab heating temperature of 1000.degree. C. or more to form a
sheet bar, finish-rolling the sheet bar so that the finisher
delivery temperature is 800.degree. C. or more, and coiling the
finish-rolled sheet at a coiling temperature of 800.degree. C. or
less to form a hot-rolled sheet; the cold rolling step of pickling
and cold-rolling the hot-rolled sheet to form a cold-rolled sheet;
and the cold-rolled sheet annealing step of continuously annealing
the cold-rolled sheet at a temperature in the ferrite-austenite
two-phase region, and cooling the annealed sheet to the temperature
region of 500.degree. C. or less at a cooling rate of 10 to
300.degree. C./s; wherein the steel slab has a composition, by mass
%, comprising at least one of: C: 0.0015 to 0.025%; Si: 1.0% or
less; Mn: 2.0% or less; P: 0.1% or less; S: 0.02% or less; Al:
0.02% or less; N: 0.0050 to 0.0250%; and at least one of the
following: B: 0.0001 to 0.0050%; and Nb: 0.002 to 0.050%; wherein
N/Al is 0.30 or more.
20. A method of producing a cold-rolled steel sheet according to
claim 19, wherein the composition further comprises, by mass %, at
least one of the following groups a to c: Group a: at least one of
Cu, Ni, Cr and Mo in a total of 1.0% or less; Group b: one or both
of Ti and V in a total of 0.1% or less; and Group c: one or both of
Ca and REM in a total of 0.0010 to 0.010%.
21. A high-tensile-strength cold-rolled steel sheet having a high r
value and excellent strain age hardenability and natural aging
resistance comprising a composition, by mass %, comprising: C:
0.025 to 0.15%; Si: 1.0% or less; Mn: 2.0% or less; P: 0.08% or
less; S: 0.02% or less; Al: 0.02% or less; and N: 0.0050 to
0.0250%; wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, the balance is composed of Fe and inevitable
impurities, the structure is composed of a ferrite phase having an
average crystal grain diameter of 10 .mu.m or less at an area ratio
of 80% or more and a martensite phase as a second phase at an area
ratio of 2% or more, and the r value is 1.2 or more.
22. A high-tensile-strength cold-rolled steel sheet according to
claim 21, wherein the composition further comprises at least one of
the following groups d to g: Group d: at least one of Cu, Ni, Cr
and Mo in a total of 1.0% or less; Group e: at least one of Nb, Ti
and V in a total of 0.1% or less; Group f: 0.0030% or less of B;
and Group g: one or both of Ca and REM in a total of 0.0010 to
0.010%.
23. A method of producing a high-tensile-strength cold-rolled steel
sheet having a r value of as high as 1.2 or more, and excellent
strain age hardenability and natural aging resistance comprising:
the hot-rolling step of roughly rolling a steel slab by heating to
a slab heating temperature of 1000.degree. C. or more to form a
sheet bar, finish-rolling the sheet bar so that: the finisher
delivery temperature is 800.degree. C. or more, and coiling the
finish-rolled sheet at a coiling temperature of 800.degree. C. or
less to form a hot-rolled sheet; the cold rolling step of pickling
and cold-rolling the hot-rolled sheet to form a cold-rolled sheet;
and the cold-rolled sheet annealing step of box-annealing the
cold-rolled sheet at an annealing temperature of the
recrystallization temperature to 800.degree. C., then continuously
annealing the annealed sheet at an annealing temperature of
Ac.sub.1 transformation point to Ac.sub.3 transformation
point--20.degree. C.), and then cooling the sheet to the
temperature region of 500.degree. C. or less at a cooling rate of
10 to 300.degree. C./s; wherein the steel slab has a composition,
by mass %, comprising at least one of: C: 0.025 to 0.15%; Si: 1.0%
or less; Mn: 2.0% or less; P: 0.08% or less; S: 0.02% or less; Al:
0.02% or less; and N: 0.0050 to 0.0250%; wherein N/Al is 0.30 or
more.
24. A method of producing a high-tensile-strength cold-rolled steel
sheet according to claim 23, further comprising performing over
aging in a temperature region of 350.degree. C. to the cooling step
temperature for a residence tine of 20 seconds or more subsequent
to cooling after the continuous annealing.
25. A method of producing a high-tensile-strength cold-rolled steel
sheet according to claim 23 or 24, wherein the composition further
comprises, by mass %, at least one of the following groups d to g:
Group d: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or
less; Group e: at least one of Nb, Ti and V in a total of 0.1% or
less; Group f: 0.0030% or less of B; and Group g: one or both of Ca
and REM in a total of 0.0010 to 0.010%.
Description
TECHNICAL FIELD
[0001] The present invention relates to a cold-rolled steel sheet,
an electro-galvanized steel sheet, a hot-dip galvanized steel
sheet, and an alloyed hot-dip galvanized steel sheet, which are
suitable as raw material steel sheets for molded products such as
building members, mechanical structural parts, automobile
structural parts, etc., which are used at positions required to
have structural strength, particularly, strength and/or stiffness
in deformation, and which are subjected to heat treatment for
increasing strength after processing such as pressing or the like,
and a method of producing these steel sheets.
[0002] 1. In the present invention, "excellent strain age
hardenability" means that in aging under conditions of holding at a
temperature of 170.degree. C. for 20 min. after pre-deformation
with a tensile strain of 5%, the increment in deformation stress
(represented by the amount of BH=yield stress after
aging-pre-deformation stress before aging) after aging is 80 MPa or
more, and the increment in tensile strength (represented by
.DELTA.TS=tensile strength after aging-tensile strength before
pre-deformation) after strain aging (pre-deformation+aging) is 40
MPa or more.
BACKGROUND ART
[0003] In producing a press-molded product of a thin steel sheet, a
process of coating and baking at lower than 200.degree. C. is used
as a method in which a material having low deformation stress
before press forming to facilitate press forming, and then hardened
after press forming to increase the strength of a part. As a steel
sheet for such coating and baking, a BH steel sheet has been
developed.
[0004] For example, Japanese Unexamined Patent Application
Publication No. 55-141526 discloses a method in which Nb is added
according to the contents of C, N and Al of steel, Nb/(dissolved
C+dissolved N) by at % is limited in a specified range, and the
cooling rate after annealing is controlled to adjust dissolved C
and dissolved N in a steel sheet. Also, Japanese Examined Patent
Application Publication No. 61-45689 discloses a method in which
baking hardenability is improved by adding Ti and Nb.
[0005] However, in order to improve deep drawability, strength of
the raw material sheets of the above-described steel sheets is
decreased, and thus the steel sheets are not always sufficient as
structural materials.
[0006] Japanese Unexamined Patent Application Publication No.
5-25549 discloses a method in which baking hardenability is
improved by adding W, Cr and Mo to steel singly or in a
combination.
[0007] In the above-described conventional techniques, strength is
increased by bake-hardening due to the functions of small amounts
of dissolved C and dissolved N in a steel sheet, and it is well
known that a BH (Bake-Hardening) steel sheet is used for increasing
only the yield strength of a material, not for increasing tensile
strength. Therefore, the conventional techniques have only the
effect of increasing the deformation start stress of a part, and
the effect of increasing stress (tensile strength after forming)
required for deformation over the entire deformation region from
the deformation start to the deformation end is not said to be
sufficient.
[0008] As a cold-rolled steel sheet having tensile strength
increased after forming, for example, Japanese Unexamined Patent
Application Publication No. 10-310847 discloses an alloying ho-dip
galvanized steel sheet having tensile strength increased by 60 MPa
or more by heat treatment in the temperature region of 200 to
450.degree. C.
[0009] This steel sheet contains, by mass %, 0.01 to 0.08% of C,
and 0.01 to 3.0% of Mn, and at least one of W, Cr, and Mo in a
total of 0.05 to 3.0%, and further contains at least one of 5 0.005
to 0.1% of Ti, 0.005 to 0.1% of Nb and 0.005 to 0.1% of V according
to demand, and the microstructure of the steel is composed of
ferrite or mainly composed of ferrite.
[0010] However, this technique comprises forming a fine carbide in
the steel sheet by heat treatment after forming to effectively
propagate a dislocation of stress applied during pressing, thereby
increasing the amount of strain. Therefore, heat treatment must be
performed in the temperature range of 220 to 370.degree. C. There
is thus the problem of a necessary heat treatment temperature
higher than general bake-hardening temperatures.
[0011] Furthermore, it is a very important problem that the body
weight of an automobile is decreased in relation to the recent
regulation of exhaust gases due to global environmental problems.
In order to decrease the body weight of an automobile, it is
effective to increase the strength of the used steel sheet, i.e.,
use a high-tensile-strength steel sheet, thinning the steel sheet
used.
[0012] An automobile part using a high-tensile-strength thin steel
sheet must exhibit a sufficient property according to its function.
The property depends upon the part, and examples of the property
include dent resistance, static strength against deformation such
as bending, twisting, or the like, fatigue resistance, impact
resistance, etc. Namely, the high-tensile-strength steel sheet used
for an automobile part is required to be excellent in such a
property after forming. The properties are related to the strength
of a steel sheet after forming, and thus the lower limit of
strength of the high-tensile-strength steel sheet used must be set
for achieving thinning.
[0013] On the other hand, in the process for forming an automobile
part, a steel sheet is press-molded. If the steel sheet has
excessively high strength in press forming, the steel sheet causes
the following problems: (1) deteriorating shape fixability; (2)
deteriorating ductility to cause cracking, necking, or the like
during forming; and (3) deteriorating dent resistance (resistance
to a dent produced by a local compressive load) when the sheet
thickness is decreased. These problems thus inhibit the extension
of application of the high-tensile-strength steel sheet to
automobile bodies.
[0014] As a means for overcoming the problems, a steel sheet
composed of ultra-low-carbon steel is known as a raw material, for
example, for a cold-rolled steel sheet for an external sheet panel,
in which the content of C finally remaining in a solid solution
state is controlled to an appropriate range. This type of steel
sheet is kept soft during press forming to ensure shape fixability
and ductility, and its yield stress is increased by utilizing the
strain aging phenomenon which occurs in the step of coating and
baking at 170.degree. C. for about 20 minutes after press forming,
to ensure dent resistance. This steel sheet is soft during press
forming because C is dissolved in steel, while dissolved C is fixed
to a dislocation introduced in press forming in the coating and
baking step after press forming to increase the yield stress.
[0015] However, in this type of steel sheet, the increase in yield
stress due to strain age hardening is kept down from the viewpoint
of prevention of the occurrence of stretcher strain causing a
surface defect. This causes the fault that the steel sheet actually
less contributes to a reduction in weight of a part.
[0016] On the other hand, a steel sheet composed of dissolved N to
further increase the amount of bake-hardening, and a steel sheet
provided with a composite structure composed of ferrite and
martensite to further improve baking hardenability have been
proposed for applications in which the appearance is not so
important.
[0017] For example, Japanese Unexamined Patent Application
Publication No. 60-52528 discloses a method of producing a
high-strength steel thin sheet having good ductility and spot
weldability, in which steel containing 0.02 to 0.15% of C, 0.8 to
3.5% of Mn, 0.02 to 0.15% of P, 0.10% or less of Al, and 0.005 to
0.025% of N is hot-rolled by coiling at a temperature of
550.degree. C. or less, cold-rolled, and then annealed by
controlled cooling and heat treatment. A steel sheet produced by
the technique disclosed in Japanese Unexamined Patent Application
Publication No. 60-52528 has a mixed structure comprising a
low-temperature transformation product phase mainly composed of
ferrite and martensite, and having excellent ductility, and high
strength is achieved by utilizing strain aging due to positively
added N during coating baking.
[0018] Although the technique disclosed in Japanese Unexamined
Patent Application Publication No. 60-52528 greatly increases yield
stress YS due to strain age hardening, the technique less increases
tensile strength TS. Also, this technique causes large variations
in the increment in yield stress YS to cause large variations in
mechanical properties, and thus it cannot be expected that a steel
sheet can be sufficiently thinned for contributing to a reduction
in weight of an automobile part, which is currently demanded.
[0019] Japanese Examined Patent Application Publication No. 5-24979
discloses a high-tensile-strength cold-rolled steel thin sheet
having baking hardenability which has a composition comprising 0.08
to 0.20% of C, 1.5 to 3.5% of Mn, and the balance composed of Fe
and inevitable impurities, and a structure composed of homogeneous
bainite containing 5% or less of ferrite, or bainite partially
containing martensite. The cold-rolled steel sheet disclosed in
Japanese Examined Patent Application Publication No. 5-24979 is
produced by quenching in the temperature range of 200 to
400.degree. C. in the cooling process after continuous annealing,
and then slowly cooling to obtain a structure mainly composed of
bainite and having a large amount of bake-hardening which is not
obtained by a conventional method.
[0020] However, in the steel sheet disclosed in Japanese Examined
Patent Application Publication No. 5-24979, yield strength is
increased after coating and baking to obtain a large amount of
bake-hardening which is not obtained a conventional method, while
tensile strength cannot be increased. Therefore, in application to
a strength member, improvements in fatigue resistance and impact
resistance after forming cannot be expected. Therefore, there is a
problem in which the steel sheet cannot be used for applications
greatly required to have fatigue resistance and impact resistance,
etc.
[0021] Also, Japanese Examined Patent Application Publication No.
61-12008 discloses a method of producing a high-tensile-strength
steel sheet having a high r value. This method is characterized by
annealing ultra-low-C steel used as a raw material in a
ferrite-austenite coexistence region after cold rolling. However,
the resultant steel sheet has a high r value and a high degree of
baking hardenability (BH property), but the obtained BH amount is
about 60 MPa at most. Also, the yield point of the steel sheet is
increased after strain aging, but TS is not increased, thereby
causing the problem of limiting application to parts.
[0022] Furthermore, the above-described steel sheet exhibits
excellent strength after coating and baking in a simple tensile
test, but produces large variations in strength during plastic
deformation under actual pressing conditions. Therefore, it cannot
be said that the steel sheet is sufficiently applied to parts
required to have reliability.
[0023] With respect of a hot-rolled steel sheet among coating baked
steel sheets for press molded products, for example, Japanese
Examined Patent Application Publication No. 8-23048 discloses a
method of producing a hot-rolled steel sheet which is soft during
processing, and has tensile strength increased by coating and
baking after processing to be effective to improve fatigue
resistance.
[0024] In this technique, steel contains 0.02 to 0.13 mass % of C,
and 0.0080 to 0.0250 mass % of N, and the finisher deliver
temperature and the coiling temperature are controlled to leave a
large amount of dissolved N in the steel, thereby forming a
composite structure as a metal structure mainly composed of ferrite
and martensite. Therefore, an increase of 100 MPa or more in
tensile strength is achieved at the heat treatment temperature of
170.degree. C. after forming.
[0025] Japanese Unexamined Patent Application Publication No.
10-183301 discloses a hot-rolled steel sheet having excellent
baking hardenability and natural aging resistance, in which the C
and N contents are limited to 0.01 to 0.12 mass % and 0.0001 to
0.01 mass %, respectively, and the average crystal grain diameter
is controlled to 8 .mu.m or less to ensure a BH amount of as high
as 80 MPa or more, and suppress the AI amount to 45 MPa or
less.
[0026] However, this steel sheet is a hot-rolled sheet, and is thus
difficult to obtain a high r value because the ferrite aggregation
texture is made random due to austeniste-ferrite transformation.
Therefore, the steel sheet cannot be said to have sufficient deep
drawability.
[0027] Furthermore, even if the hot-rolled steel sheet obtained by
this technique is used as a starting material for cold rolling and
recrystallization annealing, the increase in tensile strength
obtained after forming and heat treatment is not always equivalent
to a hot-rolled steel sheet, and a BH amount of as high as 80 MPa
or more cannot be always obtained. This is because the
microstructure of the cold-rolled steel becomes different from that
of hot-rolled one due to cold rolling and recrystallization
annealing, and strain greatly accumulates during cold rolling to
easily form a carbide, a nitride or a carbonitride, thereby
changing the states of dissolved C and dissolved N.
[0028] In consideration of the above-described present conditions,
an object of the present invention is to provide a cold-rolled
steel sheet and a hot-dip galvanized steel sheet (including an
alloyed steel sheet) for deep drawing, which have excellent deep
drawability, TS.times.r value.gtoreq.750 MPa, and excellent strain
aging hardenability (BH.gtoreq.80 MPa and .DELTA.TS.gtoreq.40 MPa),
and an effective method of producing these steel sheets.
[0029] Another object of the present invention is to solve the
above problems of the conventional techniques and provide a
high-tensile-strength cold-rolled steel sheet which is suitable for
automobile parts required to have high moldability, softness and
high moldability, and stable material properties, and which can
easily be molded to an automobile part having a complicated shape
without producing shape defects such as spring back, twisting, and
curving, and cracking, etc., and which has sufficient strength as
an automobile part after heat treatment of a molded automobile part
to permit sufficient contribution to a reduction in body weight of
an automobile, a high r value of 1.2 or more, and excellent strain
age hardenability. A further object of the present invention is to
provide an industrial production method capable of producing the
steel sheet at low cost without disturbing its shape.
DISCLOSURE OF INVENTION
[0030] In order to achieve the objects, the inventors produced
various steel sheets having different compositions under various
production conditions, and experimentally evaluated various
material properties. As a result, it was found that both
moldability and hardenability after forming can be improved by
using as a strengthening element N, which has not be positively
used before in a field requiring high processability, and
effectively using the great strain age hardening phenomenon
manifested by the action of the strengthening element.
[0031] The inventors also found that in order to advantageously use
the strain age hardening phenomenon due to N, the strain age
hardening phenomenon due to N must be advantageously combined with
a condition for coating and baking an automobile, or further
positively combined with a heat treatment condition after forming.
It was thus found to be effective to appropriately control the hot
rolling condition, the cold rolling and the cold rolling annealing
condition to control the microstructure of a steel sheet and the
amount of dissolved N in certain ranges. It was also found that in
order to stably manifest the strain age hardening phenomenon due to
N, it is important to control the Al content of the composition
according to the N content.
[0032] The inventors further found that in order to obtain a high r
value, the C content is decreased, continuous annealing is
performed in the ferrite-austenite two-phase temperature region,
and subsequent cooling is controlled to form a structure containing
an acicular ferrite phase at an area ratio of 5% or more in the
ferrite phase. Such a combination of the microstructure and the
appropriate amount of dissolved N was found to enable the
achievement of a cold-rolled steel sheet having a high r value,
excellent press moldability, and excellent strain age
hardenability. This was also found to permit sufficient use of N
without causing the problem of natural aging deterioration, which
is the problem of a conventional bake-hardening steel sheet.
[0033] Namely, the inventors found that by suing N as a
strengthening element, controlling the Al content according to the
N content in an appropriate range, and appropriately controlling
the hot rolling condition and the cold rolling annealing condition
to appropriately control the microstructure and dissolved N, it is
possible to obtain a steel sheet having a high r value and
excellent moldability as compared with conventional solid-solution
strengthening-type C--Mn steel sheets and precipitation
strengthening-type steel sheets, and strain age hardenability,
which is not possessed by the conventional steel sheets.
[0034] A steel sheet of the present invention exhibits higher
strength after coating and baking in a simple tensile test, as
compared with a conventional steel sheet, and exhibits small
variations in strength in plastic deformation under actual pressing
conditions and stable part strength, thereby enabling application
to parts required to have reliability. For example, a portion where
large strain is applied to decrease the thickness has higher
hardenability than other portions, and is considered homogeneous
when being evaluated based on a surcharge load ability of
(thickness).times.(strength), thereby stabilizing strength as a
part.
[0035] As a result of further intensive research for achieving the
objects, the inventors found the following:
[0036] 1) In order to increase tensile strength after forming and
heat treatment, a new dislocation must be introduced for
progressing tensile deformation. The movement of the dislocation
introduced by pre-deformation must be prevented by interaction
between the dislocation introduced by forming and an interstitial
element or a precipitate even when upper yield stress is
attained.
[0037] 2) In order to obtain the above interaction by forming a
carbide, a nitride or a carbonitride of W, Cr, Mo, Ti, Nb, Al or
the like, the heat treatment temperature after forming must be
increased to 200.degree. C. or more. Therefore, it is more
advantageous to positively use the interstitial element or a Fe
carbide or Fe nitride because the heat treatment temperature after
forming is decreased.
[0038] 3) Of interstitial elements, dissolved N has the higher
interaction with a dislocation introduced by forming than dissolved
C even when the heat treatment temperature after forming is
decreased, and thus a dislocation introduced by pre-deformation
less moves when upper yield stress is attained.
[0039] 4) Although dissolved N is present in crystal grains and
crystal grain boundaries in steel, the increase in strength after
forming and heat treatment increases as the area of the crystal
grain boundaries increases. Namely, the smaller crystal grain
diameter is advantageous.
[0040] 5) In order to increase the crystal grain boundary area, it
is advantageous to add a combination of Nb and B and cool
immediately after the end of hot rolling, suppressing normal grain
growth of ferrite grains after the end of hot rolling and
suppressing grain growth by recrystallization annealing after cold
rolling.
[0041] The present invention has been achieved based on the above
findings. The findings were obtained from the experiment described
below.
[0042] Experiment 1
[0043] A sheet bar (thickness: 30 mm) having a composition
containing, by mass %, 0.0015% of C, 0.0010% of B, 0.015 of Si,
0.5% of Mn, 0.03% of P, 0.08% of S and 0.011% of N, 0.005 to 0.05%
of Nb and 0.005 to 0.03% of Al, and the balance composed of Fe and
inevitable impurities was uniformly heated at 1150.degree. C.,
hot-rolled by three passes so that the temperature of the final
pass was 900.degree. C. higher than the Ar.sub.3 transformation
point, and then cooled with water for 0.1 second. Then, the sheet
bar was subjected to heat treatment corresponding to coiling at
500.degree. C. for 1 hour.
[0044] The thus-obtained hot-rolled sheet having a thickness of 4
mm was cold-rolled with a rolling reduction ratio of 82,5%,
recrystallized and annealed at 800.degree. C. for 40 seconds, and
then temper-rolled with a rolling reduction ratio of 0.8%. Then, a
tensile test specimen of JIS No. 5 was obtained from the resultant
cold-rolled sheet in the rolling direction, and tensile strength
was measured with a strain rate of 0.02/s by using a general
tensile testing machine. Also, tensile strain of 10% was applied to
a tensile test specimen of JIS No. 5 separately obtained from the
cold-rolled sheet in the rolling direction, and then the specimen
was subjected to a normal tensile test after heat treatment at
120.degree. C. for 20 minutes. The difference between the tensile
strength of the specimen obtained from the cold-rolled sheet and
the tensile strength of the specimen heat treated at 120.degree. C.
for 20 minutes after application of 10% tensile strain was
considered as the increase in tensile strength after forming
(.DELTA.TS).
[0045] FIG. 1 shows the results of measurement of the relation
between the steel compositions (N %-14/93.multidot.Nb
%-14/27.multidot.Al %-14/11.multidot.B %) and .DELTA.TS.
[0046] The figure indicates that .DELTA.TS becomes 60 MPa or more
when the value of (N % -14/93.multidot.Nb %-14/27.multidot.Al %
-14/11.multidot.B %). satisfies 0.0015 mass %.
[0047] Experiment 2
[0048] A sheet bar (thickness: 30 mm) having a composition
containing, by mass %, 0.0010% of C, 0.02 of Si, 0.6% of Mn, 0.01%
of P, 0.009% of S and 0.012% of N, 0.01% of Al, 0.015% of Nb,
0.00005 to 0.0025% of B, and the balance composed of Fe and
inevitable impurities was uniformly heated at 1100.degree. C.,
hot-rolled by three passes so that the temperature of the final
pass was 920.degree. C. higher than the Ar.sub.3 transformation
point, and then cooled with water for 0.1 second. Then, the sheet
bar was subjected to heat treatment corresponding to coiling at
450.degree. C. for 1 hour.
[0049] The thus-obtained hot-rolled sheet having a thickness of 4
mm was cold-rolled with a rolling reduction ratio of 82,5%,
recrystallized and annealed at 820.degree. C. for 40 seconds, and
then temper-rolled with a rolling reduction ratio of 0.8%. Then, a
tensile test specimen of JIS No. 5 was obtained from the resultant
cold-rolled sheet in the rolling direction, and tensile strength
was measured with a strain rate of 0.02/s by using a general
tensile testing machine. Also, tensile strain of 10% was applied to
a tensile test specimen of JIS No. 5 separately obtained from the
cold-rolled sheet in the rolling direction, and then the specimen
was subjected to a normal tensile test after heat treatment at
120.degree. C. for 20 minutes.
[0050] FIG. 2 shows the results of measurement of the relation
between the B content of steel and .DELTA.TS. This figure indicates
that with a B content of 0.0005 to 0.0015 mass %, a high .DELTA.TS
of 60 MPa or more can be obtained.
[0051] As a result of observation of the microstructure, it was
also found that by adding a combination of Nb and B to make fine
crystal grains, a high .DELTA.TS can be obtained.
[0052] Namely, with a B content of less than 0.0005 mass %, the
effect of making fine crystal grains by adding a combination with
Nb is small. On the other hand, with a B content of over 0.0015
mass %, the amount of B segregated in the grain boundaries and the
vicinities thereof is increased to decrease the amount of effective
dissolved N because of the strong interaction between B atoms and N
atoms, thereby possibly decreasing .DELTA.TS.
[0053] Experiment 3
[0054] A sheet bar (thickness: 30 mm) of each of steel A having a
composition containing, by mass %, 0.0010% of C, 0.012% of N,
0.0010% of B, 0.01% of Si, 0.5% of Mn, 0.03% of P, 0.008% of S,
0.014% of Nb, 0.01% of Al, and the balance composed of Fe and
inevitable impurities, and steel B having a composition containing,
by mass %, 0.010% of C, 0.0012% of N, 0.0010% of B, 0.01% of Si,
0.5% of Mn, 0.03% of P, 0.008% of S, 0.014% of Nb, 0.01% of Al, and
the balance composed of Fe and inevitable impurities was uniformly
heated at 1150.degree. C., hot-rolled by three passes so that the
temperature of the final pass was 910.degree. C. higher than the
Ar.sub.3 transformation point, and then cooled with a gas for 0.1
second. Then, each of the sheet bars was subjected to heat
treatment corresponding to coiling at 600.degree. C. for 1
hour.
[0055] Each of the thus-obtained hot-rolled sheets having a
thickness of 4 mm was cold-rolled with a rolling reduction ratio of
82,5%, recrystallized and annealed at 880.degree. C. for 40
seconds, and then temper-rolled with a rolling reduction ratio of
0.8%.
[0056] Then, a tensile test specimen of JIS No. 5 was obtained from
each of the resultant cold-rolled sheets in the rolling direction,
and tensile strength was measured with a strain rate of 0.02/s by
using a general tensile testing machine. Also, tensile strain of
10% was applied to a tensile test specimen of JIS No. 5 separately
obtained from each of the cold-rolled sheets in the rolling
direction, and then the specimen was subjected to a normal tensile
test after heat treatment at various temperatures for 20
minutes.
[0057] FIG. 3 shows the results of measurement of the influence of
the heat treatment temperature after forming on .DELTA.TS. This
figure indicates that in the relatively low temperature region of
heat treatment temperatures of 200.degree. C. or less after
forming, the ultra-low carbon steel A having a high N content
exhibits higher .DELTA.TS than the semi-ultra low carbon steel B
having a low N content, and while in the high temperature region,
both steel materials exhibit substantially the same .DELTA.TS.
There experimental results reveal that in order to ensure .DELTA.TS
in the low temperature region, it is effective to use dissolved
N.
[0058] FIG. 4 shows the results of measurement of the influences of
the crystal grain diameter d and steel compositions (N
%-14/93.multidot.Nb %-14/27.multidot.Al %-14/11.multidot.B %) on a
decrease (.DELTA.El) in elongation by natural aging and an increase
in tensile strength (.DELTA.TS) after forming. The decrease
(.DELTA.El) in elongation was evaluated by the difference between
the total elongation measured with the test specimen of JIS NO. 5
obtained from each of the cold-rolled sheets in the rolling
direction, and the total elongation measured with the separately
obtained test specimen after holding at 100.degree. C. for 8 hours
for accelerating natural aging.
[0059] FIG. 4 indicates that when the value of (N
%-14/93.multidot.Nb % -14/27.multidot.Al %-14/11.multidot.B %) is
0.0015 mass % or more, and the crystal grain diameter d is 20 .mu.m
or less, both high .mu.TS and low .DELTA.El can be achieved.
[0060] Experiment 4
[0061] A sheet bar of steel containing 0.0015% of C, 0.30 of Si,
0.8% of Mn, 0.03% of P, 0.005% of S and 0.012% of N, and 0.02 to
0.08% of Al was uniformly heated at 1050.degree. C., hot-rolled by
seven passes so that the temperature of the final pass was
670.degree. C., and then recrystallized and annealed at 700.degree.
C. for 5 hours. The thus-obtained hot-rolled sheet having a
thickness of 4 mm was cold-rolled with a rolling reduction ratio of
82,5%, recrystallized and annealed at 875.degree. C. for 40
seconds, and then temper-rolled with a rolling reduction of 0.8%.
Then, a tensile test specimen of JIS No. 5 was obtained from the
resultant cold-rolled sheet in the rolling direction, and
TS.times.r value and .DELTA.TS were measured with a strain rate of
3.times.10.sup.-3/s by using a general tensile testing machine. The
results are shown in FIG. 5. In this figure, when N/Al.gtoreq.0.03
is satisfied, TS.times.r value.gtoreq.750 and .DELTA.TS.gtoreq.40
MPa are achieved. It was also confirmed that when N/Al.gtoreq.0.03,
BH.gtoreq.80 MPa is attained.
[0062] Experiment 5
[0063] A sheet bar of steel containing 0.0015% of C, 0.0010% of B,
0.01 of Si, 0.5% of Mn, 0.03% of P, 0.008% of S and 0.011% of N,
0.005 to 0.05% of Nb, and 0.005 to 0.03% of Al was uniformly heated
at 1000.degree. C., hot-rolled by seven passes so that the
temperature of the final pass was 650.degree. C., and then
recrystallized and annealed at 800.degree. C. for 60 seconds. The
thus-obtained hot-rolled sheet having a thickness of 4 mm was
cold-rolled with a rolling reduction ratio of 82,5%, recrystallized
and annealing at 880.degree. C. for 40 seconds, and then
temper-rolled with a rolling reduction ratio of 0.8%. Then, a
tensile test specimen of JIS No. 5 was obtained from the resultant
cold-rolled sheet in the rolling direction, and TS.times.r value,
BH and .DELTA.TS were measured with a strain rate of
3.times.10.sup.-3/s by using a general tensile testing machine. The
relations between the measured values and N/(Al+Nb+B) are shown in
FIG. 5. In this experiment, steel containing 0.005 to 0.05% of Nb
and 0.0010% of B was used. This figure indicates that in the range
of N/(Al+Nb+B).gtoreq.0.30, BH.gtoreq.80 MPa, .DELTA.TS.gtoreq.60
MPa, and TS.times.r value.gtoreq.850 are achieved.
[0064] Experiment 6
[0065] A sheet bar of steel containing 0.0010% of C, 0.02 of Si,
0.6% of Mn, 0.01% of P, 0.009% of S and 0.015% of N, 0.01% of Al,
0.015% of Nb and 0.0001 to 0.0025% of B was uniformly heated at
1050.degree. C., hot-rolled by seven passes so that the temperature
of the final pass was 680.degree. C., and then recrystallized and
annealed at 850.degree. C. for 5 hours. The thus-obtained
hot-rolled sheet having a thickness of 4 mm was cold-rolled with a
rolling reduction ratio of 82,5%, recrystallized and annealed at
880.degree. C. for 40 seconds, and then temper-rolled with a
rolling reduction ratio of 0.8%. Then, a tensile test specimen of
JIS No. 5 was obtained from the resultant cold-rolled sheet in the
rolling direction, and TS.times.r value, BH and .DELTA.TS were
measured with a strain rate of 3.times.10.sup.-3/s by using a
general tensile testing machine. The relations between the measured
values and the B content are shown in FIG. 6.
[0066] This figure indicates that in the B content range of 0.0003
to 0.0015%, BH.gtoreq.80 MPa, .DELTA.TS.gtoreq.60 MPa, which is
higher than the case of B<0.0003%, and TS.times.r
value.gtoreq.850 are achieved. As a result of observation of the
microstructure, it was also confirmed that in this B range, crystal
grains are significantly made fine.
[0067] The results shown in FIGS. 5 and 6 indicate that in the
range of N/(Al+Nb+B).gtoreq.0.30 wherein B.gtoreq.0.0003%, the
crystal grains are further made fine by combining Nb, and .DELTA.TS
and the level of TS.times.r value are further improved. When
B<0.0003%, the effect of making fine crystal grains by combining
Nb is not exhibited. On the other hand, when B>0.0015%,
properties further deteriorate. This is possibly due to the fact
that the amount of B segregated in the grain boundaries and the
vicinities thereof is increased to decrease the amount of effective
dissolved N due to the strong interaction between B and N atoms.
The same research as described above was carried out for the case
in which Ti and V were added in place of Nb, and it was confirmed
that the same effect as Nb could be obtained. The present invention
has been achieved based on the above-described findings, and the
gist of the invention was follows.
[0068] In a first aspect of the present invention, a cold-rolled
steel sheet having excellent strain age hardenability comprises a
composition, by mass %, comprising:
[0069] C: 0.15% or less;
[0070] Si: 1.0% or less;
[0071] Mn: 2.0% or less;
[0072] P: 0.1% or less;
[0073] S: 0.01% or less;
[0074] Al: 0.005 to 0.030%; and
[0075] N: 0.0050 to 0.0400%;
[0076] wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, and the balance is composed of Fe and inevitable
impurities.
[0077] In the first aspect of the present invention, the
above-described composition preferably comprises the compositions
in the following ranges. Namely, a cold-rolled steel sheet having
excellent strain age hardenability comprises a composition, by mass
%, comprising:
[0078] C: less than 0.01%;
[0079] Si: 0.005 to 1.0%;
[0080] Mn: 0.01 to 1.5%;
[0081] P: 0.1% or less;
[0082] S: 0.01% or less;
[0083] Al: 0.005 to 0.030%; and
[0084] N: 0.005 to 0.040%;
[0085] wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, and the balance is composed of Fe and inevitable
impurities.
[0086] In the first aspect of the present invention, the
composition, by mass %, preferably further comprises:
[0087] B: 0.0001 to 0.0030%; and
[0088] Nb: 0.005 to 0.050%;
[0089] wherein the ranges of B and Nb satisfy the following
equations (1) and (2):
N %.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1)
C %.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
[0090] In the first aspect of the present invention, the above
composition, by mass %, preferably further comprises at least one
of Cu, Ni and Mo in a total of 1.0% or less according to
demand.
[0091] In the first aspect of the present invention, the steel
sheet preferably has a crystal grain diameter of 20 .mu.m or
less.
[0092] In the first aspect of the present invention, strength after
forming is preferably increased by 60 MPa or more by heat treatment
in the low temperature region of 120 to 200.degree. C.
[0093] In the first aspect of the present invention, the surface of
the cold-rolled steel sheet may be coated by electro-galvanization,
hot-dip galvanization, or alloying hot-dip galvanization.
[0094] In a second aspect of the present invention, a method of
producing a cold-rolled steel sheet having excellent strain age
hardenability comprises hot-rolling a steel slab under conditions
in which the slab is cooled immediately after the end of finish
rolling and coiled at a coiling temperature of 400 to 800.degree.
C., cold-rolling the hot-rolled sheet with a rolling reduction
ratio of 60 to 95%, and then performing recrystallization annealing
at a temperature of 650 to 900.degree. C., wherein the steel slab
has a composition, by mass %, comprising:
[0095] C: less than 0.01%;
[0096] Si: 0.005 to 1.0%;
[0097] Mn: 0.01 to 1.5%;
[0098] P: 0.1% or less;
[0099] S: 0.01% or less;
[0100] Al: 0.005 to 0.030%; and
[0101] N: 0.005 to 0.040%;
[0102] wherein N/Al is 0.30 or more, and the balance is
substantially composed of Fe.
[0103] In the second aspect of the present invention, the
composition, by mass %, preferably further comprises:
[0104] B: 0.0001 to 0.0030%; and
[0105] Nb: 0.005 to 0.050%;
[0106] wherein the ranges of B and Nb satisfy the following
equations (1) and (2):
N %.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1)
C %.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
[0107] In the second aspect of the present invention, in the
heating-up step in recrystallization annealing, the temperature is
preferably increased at a rate of 1 to 20.degree. C./s in the
temperature region from 500.degree. C. to the recrystallization
temperature.
[0108] In the second aspect of the present invention, hot-dip
galvanization and heat alloying may be performed after the
recrystallization annealing.
[0109] In a third aspect of the present invention, a cold-rolled
deep drawing steel sheet having excellent strain age hardenability
comprises a composition, by mass %, comprising:
[0110] C: 0.01% or less;
[0111] Si: 1.0% or less;
[0112] Mn: 0.01 to 1.5%;
[0113] P: 0.1% or less;
[0114] S: 0.01% or less;
[0115] Al: 0.005 to 0.020%; and
[0116] N: 0.0050 to 0.040%;
[0117] wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, the balance is composed of Fe and inevitable
impurities, and Ts x r value is 750 MPa or more.
[0118] In the third aspect of the present invention, the
composition, by mass %, preferably further comprises:
[0119] B: 0.0001 to 0.0030%; and
[0120] Nb: 0.005 to 0.050%;
[0121] wherein the ranges of B and Nb satisfy the following
equations (1) and (2):
N %.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1)
C%.ltoreq.0.5.about.(12/93).multidot.Nb % (2)
[0122] In the third aspect of the present invention, the
composition, by mass %, preferably further comprises at least one
of the following:
[0123] B: 0.0001 to 0.0030%;
[0124] Nb: 0.005 to 0.050%;
[0125] Ti: 0.005 to 0.070%; and
[0126] V: 0.005 to 0.10%;
[0127] wherein N/(Al+Nb+Ti+V+B) is 0.30 or more, and the amount of
dissolved N is 0.0010% or more.
[0128] In a fourth aspect of the present invention, a method of
producing a cold-rolled deep drawing steel sheet having excellent
strain age hardenability comprises heating a steel raw material to
950.degree. C. or more, roughly rolling the raw material so that
the finisher delivery temperature is Ar.sub.3 to 1000.degree. C.,
finish-rolling the material while lubricating it in the temperature
region of 600.degree. C. to Ar.sub.3, coiling the rolled sheet in
which the total reduction ratio by rolling starting from rough
rolling to finish rolling is 80% or more, recrystallizing and
annealing the hot-rolled sheet, cold-rolling the rolled sheet with
a rolling reduction ratio of 60 to 95%, and then recrystallizing
and annealing the resultant cold-rolled sheet, wherein the steel
raw material has a composition, by mass %, comprising:
[0129] C: 0.01% or less;
[0130] Si: 0.005 to 1.0%;
[0131] Mn: 0.01 to 1.0%;
[0132] P: 0.1% or less;
[0133] S: 0.01% or less;
[0134] Al: 0.005 to 0.030%;
[0135] N: 0.005 to 0.040%; and
[0136] at least one of the following:
[0137] B: 0.0003 to 0.0030%;
[0138] Nb: 0.005 to 0.050%;
[0139] Ti: 0.005 to 0.070%; and
[0140] V: 0.005 to 0.10%;
[0141] wherein N/(Al+Nb+Ti+V+B) is 0.30 or more.
[0142] In a fifth aspect of the present invention, a
high-tensile-strength cold-rolled steel sheet having excellent
moldability, strain age hardenability and natural aging resistance
comprises a composition, by mass %, comprising:
[0143] C: 0.0015 to 0.025%;
[0144] Si: 1.0% or less;
[0145] Mn: 2.0% or less;
[0146] P: 0.1% or less;
[0147] S: 0.02% or less;
[0148] Al: 0.02% or less;
[0149] N: 0.0050 to 0.0250%; and
[0150] one or both of the following:
[0151] B: 0.0005 to 0.0050%; and
[0152] Nb: 0.002 to 0.050%;
[0153] wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, the balance is composed of Fe and inevitable
impurities, the structure is composed of an acicular ferrite phase
at an area ratio of 5% or more and a ferrite phase having an
average crystal grain diameter of 20 .mu.m or less, and the r value
is 1.2 or more.
[0154] In the fifth aspect of the present invention, the
composition preferably further comprises at least one of the
following groups a to c:
[0155] Group a: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0156] Group b: one or both of Ti and V in a total of 0.1% or less;
and
[0157] Group c: one or both of Ca and REM in a total of 0.0010 5 to
0.010%.
[0158] In a sixth aspect of the present invention, a method of
producing a high-tensile-strength cold-rolled steel sheet:
[0159] having a r value of 1.2 or more, and excellent moldability,
strain age hardenability and natural aging resistance
comprises:
[0160] the hot-rolling step of roughly rolling a steel slab by
heating to a slab heating temperature of 1000.degree. C. or more to
form a sheet bar, finish-rolling the sheet bar so that the finisher
delivery temperature is 800.degree. C. or more, and coiling the
finish-rolled sheet at a coiling temperature of 650.degree. C. or
less to form a hot-rolled sheet;
[0161] the cold rolling step of pickling and cold-rolling the
hot-rolled sheet to form a cold-rolled sheet; and the cold-rolled
sheet annealing step of continuously annealing the cold-rolled
sheet at a temperature in the ferrite-austenite two-phase region,
and cooling the annealed sheet to the temperature region of
500.degree. C. or less at a cooling rate of 10 to 300.degree.
C./s;
[0162] wherein the steel slab has a composition, by mass %,
comprising at least one of:
[0163] C: 0.0015 to 0.025%;
[0164] Si: 1.0% or less;
[0165] Mn: 2.0% or less;
[0166] P: 0.1% or less;
[0167] S: 0.02% or less;
[0168] Al: 0.02% or less;
[0169] N: 0.0050 to 0.0250%;
[0170] B: 0.0003 to 0.0050%; and
[0171] Nb: 0.002 to 0.050%;
[0172] wherein N/Al is 0.30 or more.
[0173] In the sixth aspect of the present invention, the
composition preferably further comprises, by mass %, at least one
of the following groups a to c:
[0174] Group a: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0175] Group b: one or both of Ti and V in a total of 0.1% or less;
and
[0176] Group c: one or both of Ca and REM in a total of 0.0010 to
0.010%.
[0177] In a seventh aspect of the present invention, a
high-tensile-strength cold-rolled steel sheet having a high r value
and excellent strain age hardenability and natural aging resistance
comprises a composition, by mass %, comprising:
[0178] C: 0.025 to 0.15%;
[0179] Si: 1.0% or less;
[0180] Mn: 2.0% or less;
[0181] P: 0.08% or less;
[0182] S: 0.02% or less;
[0183] Al: 0.02% or less; and
[0184] N: 0.0050 to 0.0250%;
[0185] wherein N/Al is 0.30 or more, the amount of dissolved N is
0.0010% or more, the balance is composed of Fe and inevitable
impurities, the structure is composed of a ferrite phase having an
average crystal grain diameter of 10 .mu.m or less at an area ratio
of 80% or more and a martensite phase as a second phase at an area
ratio of 2% or more, and the r value is 1.2 or more.
[0186] In the seventh aspect of the present invention, the
composition preferably further comprises at least one of the
following groups d to g:
[0187] Group d: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0188] Group e: at least one of Nb, Ti and V in a total of 0.1% or
less;
[0189] Group f: 0.0030% or less of B; and
[0190] Group g: one or both of Ca and REM in a total of 0.0010 to
0.010%.
[0191] In an eighth aspect of the present invention, a method of
producing a high-tensile-strength cold-rolled steel sheet having a
r value of as high as 1.2 or more, and excellent strain age
hardenability and natural aging resistance comprises:
[0192] the hot-rolling step of roughly rolling a steel slab by
heating to a slab heating temperature of 1000.degree. C. or more to
form a sheet bar, finish-rolling the sheet bar so that the finisher
delivery temperature is 800.degree. C. or more, and coiling the
finish-rolled sheet at a coiling temperature of 650.degree. C. or
less to form a hot-rolled sheet;
[0193] the cold rolling step of pickling and cold-rolling the
hot-rolled sheet to form a cold-rolled sheet; and
[0194] the cold-rolled sheet annealing step of box-annealing the
cold-rolled sheet at an annealing temperature of the
recrystallization temperature to 800.degree. C., then continuously
annealing the annealed sheet at an annealing temperature of
Ac.sub.1 transformation point to (AC.sub.3 transformation
point--20.degree. C.), and then cooling the sheet to the
temperature region of 500.degree. C. or less at a cooling rate of
10 to 300.degree. C./s;
[0195] wherein the steel slab has a composition, by mass %,
comprising at least one of:
[0196] C: 0.025 to 0.15%;
[0197] Si: 1.0% or less;
[0198] Mn: 2.0% or less;
[0199] P 0.08% or less;
[0200] S: 0.02% or less;
[0201] Al: 0.02% or less; and
[0202] N: 0.0050 to 0.0250%;
[0203] wherein N/Al is 0.30 or more.
[0204] In the eighth aspect of the present invention, over aging
after the cooling step of the continuous annealing is preferably
carried out in the temperature range of 350.degree. C. or more for
a residual time of 20 seconds or more.
[0205] In the eighth aspect of the present invention, the
composition preferably further comprises, by mass %, at least one
of the following groups d to g:
[0206] Group d: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0207] Group e: at least one of Nb, Ti and V in a total of 0.1% or
less;
[0208] Group f: 0.0030% or less of B; and
[0209] Group g: one or both of Ca and REM in a total of 0.0010 to
0.010%.
BRIEF DESCRIPTION OF THE DRAWINGS
[0210] FIG. 1 shows the relation between steel compositions (N
%-14/93.about.Nb %-14/27.multidot.Al %-14/11.multidot.B %) and the
increase in tensile strength (.DELTA.TS) after forming.
[0211] FIG. 2 shows the relation between the B content and
.DELTA.TS of steel containing a combination Nb and B.
[0212] FIG. 3 shows comparison of the difference in increase in
tensile strength by heat treatment after forming in a low
temperature region between steel B (conventional steel) containing
a large amount of dissolved C and steel A (steel of this invention)
containing a large amount of dissolved N.
[0213] FIG. 4 shows the influence of the crystal grain diameter d
and steel compositions (N %-14/93.multidot.Nb %-14/27.multidot.Al
%-14/11.multidot.B %) on the decrease in elongation (.DELTA.El) due
to natural aging and the increase in tensile strength (.DELTA.TS)
after forming.
[0214] FIG. 5 shows the relations between Ts.times.r value, BH,
.DELTA.TS and N/(Al+Nb+B).
[0215] FIG. 6 shows the relations between Ts x r value, BH,
.DELTA.TS and the B amount.
BEST MODE FOR CARRYING OUT THE INVENTION
[0216] Description will now be made of the reasons for limiting
compositions to the ranges below in accordance with a first
embodiment of the present invention.
[0217] C: Less Than 0.01 Mass %
[0218] From the viewpoint of excellent deep drawability and press
moldability, C is advantageously as small as possible. Also,
redissolution of NbC proceeds in the annealing step after cold
rolling to increase the amount of dissolved C in crystal grains,
thereby easily causing deterioration in natural aging resistance.
Therefore, the C amount is preferably suppressed to less than 0.01
mass %, more preferably 0.0050 mass % or less, and most preferably
0.0030 mass % or less.
[0219] Si: 0.005 to 1.0 Mass %
[0220] Si is a useful composition for suppressing a decrease in
elongation, and improving strength. However, with a Si content of
less than 0.005 mass %, the effect of addition of Si is
insufficient, while with a Si content of over 1.0 mass %, surface
properties deteriorate to deteriorate ductility. Therefore, the Si
content is limited to the range of 0.005 to 1.0 mass %, and
preferably the range of 0.01 to 0.75 mass %.
[0221] Mn: 0.01 to 1.5 Mass %
[0222] Mn not only is useful as a strengthening composition for
steel, but also has the function to suppress embrittlement with S
due to the formation of MnS. However, with a Mn content of less
than 0.01 mass %, the effect of addition of Mn is insufficient,
while with a Mn content of over 1.5 mass %, surface properties
deteriorate to deteriorate ductility. Therefore, the Mn content is
limited to the range of 0.01 to 1.5 mass %, and preferably the
range of 1.10 to 0.75 mass %.
[0223] P: 0.01 Mass % or Less
[0224] P is a solid solution strengthening element which
effectively contributes to reinforcement of steel. However, with a
P content of over 0.01 mass %, deep drawability deteriorates due to
the formation of phosphide such as (FeNb).sub.xP or the like.
Therefore, P is limited to 0.10 mass % or less.
[0225] S: 0.01 Mass % or Less
[0226] With a high S content, the amount of inclusions is increased
to deteriorate ductility. Therefore, contamination with S is
preferably prevented as much as possible, but an S content up to
0.01 mass % is allowable.
[0227] Al: 0.005 to 0.030 Mass %
[0228] Al is added as a deoxidizer for improving the yield of
carbonitride forming components. However, with an Al content of
less than 0.005 mass %, the effect is insufficient, while with an
Al content of over 0.030 mass %, the amount of N to be added to
steel is increased to easily cause slab defects during steel
making. Therefore, Al is contained in the range of 0.005 to 0.030
mass %.
[0229] N: 0.005 to 0.040 Mass %
[0230] In the present invention, N is an important element which
plays the role of imparting strain age hardenability to a steel
sheet. However, with an N content of less than 0.005 mass %, a
sufficient strain age hardenability cannot be obtained, while with
an N content of as high as over 0.040 mass %, press moldability
deteriorates. Therefore, N is contained in the range of 0.005 to
0.040 mass %, and preferably in the range of 0.008 to 0.015 mass
%.
[0231] B: 0.0001 to 0.003 Mass %
[0232] B is added in a combination with Nb to exhibit the function
to effectively make fine the hot-rolled structure and the
cold-rolled recrystallized annealed structure and to improve
cold-work embrittlement resistance. However, with a B content of
less than 0.0001 mass %, the sufficient effect of making fine the
structures cannot be obtained, while with a B content of over 0.003
mass %, the amount of BN precipitate is increased, and dissolution
in the slab heating step is hindered. Therefore, B is contained in
the range of 0.0001 to 0.003 mass %, preferably in the range of
0.0001 to 0.0015 mass %, and more preferably in the range of 0.0007
to 0.0012 mass
[0233] Nb: 0.005 to 0.050 Mass %
[0234] Nb is added in a combination with B to contribute to
refinement of the hot-rolled structure and the cold-rolled
recrystallized annealed structure, and have the function to fix
dissolved C as NbC. Furthermore, Nb forms a nitride NbN to
contribute to refinement of the cold-rolled recrystallized annealed
structure. However, with a Nb content of less than 0.005 mass %,
not only it becomes difficult to precipitate and fix dissolved C,
but also the hot-rolled structure and the cold-rolled,
recrystallized, annealed structure are not sufficiently made fine,
while with a Nb content of over 0.050 mass %, ductility
deteriorates. Therefore, Nb is contained in the range of 0.005 to
0.050 mass %, and preferably 0.010 to 0.030 mass %.
[0235] As described above, Nb has the function to fix dissolved C
as NbC, and forms a nitride NbN. Similarly, Al and B form AlN and
BN, respectively. Therefore, in order to ensure the sufficient
amount of dissolved N and sufficiently decrease the amount of
dissolved C, it is important to satisfy the following relations (1)
and (2):
N %.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1)
C %.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
[0236] In the present invention, in order to obtain a high strain
aging property and prevent aging deterioration, the crystal grain
diameter is preferably decreased.
[0237] Namely, as described above with reference to FIG. 4, even
when (N %-14/93.multidot.Nb %-14/27.multidot.Al %-14/11.multidot.B
%).gtoreq.0.0015 mass %, i.e., when a relatively large amount of
dissolved N is contained, AE1 can be suppressed to 2.0% or less by
decreasing the crystal grain diameter d to 20 .mu.m or less. The
crystal grain diameter d is more preferably decreased to 15 .mu.m
or less. This is because, as shown in FIG. 4, .DELTA.El can be
suppressed to 2.0% or less by decreasing the crystal grain diameter
d to 15 .mu.m or less.
[0238] The production conditions according to a second embodiment
of the present invention will be described.
[0239] Steel having the above-described suitable composition is
melted by a known melting method such as a converter or the like,
and a steel slab is formed by an ingot making method or a
continuous casting method.
[0240] Then, the steel slab is heated and soaked, and then
hot-rolled to form a hot-rolled sheet. In the present invention,
the heating temperature of hot rolling is not specified, but the
heating temperature of hot rolling is preferably set to
1300.degree. C. or less. This is because it is advantageous to fix
and precipitate dissolved C as a carbide in order to improve deep
drawability. In order to further improve processability, the
heating temperature is preferably set to 1150.degree. C. or less.
However, with a heating temperature of less than 900.degree. C.,
improvement in processability is saturated to conversely increase
the rolling load in hot rolling, thereby increasing the danger of
causing a rolling trouble. Therefore, the lower limit of the
heating temperature is preferably 900.degree. C.
[0241] The total rolling reduction ratio of hot rolling is
preferably 70% or more. This is because with a toal rolling
reduction ratio of less than 70%, the crystal grains of the
hot-rolled sheet are not sufficiently made fine.
[0242] During hot rolling, finish rolling is preferably finished in
the temperature region of 650 to 960.degree. C., and the finishing
temperature of hot-rolling may be in the .gamma. region above the
Ar.sub.3 transformation point, or the a region below the Ar.sub.3
transformation pint. With the finishing temperature in hot-rolling
process over 960.degree. C., the crystal grains of the hot-rolled
sheet are coarsened to deteriorate deep drawability after cold
rolling and annealing. On the other hand, with a temperature of
less than 650.degree. C., deformation resistance is increased to
increase the hot-rolling load, causing difficulties in rolling.
[0243] Preferably, cooling is started immediately after the end of
final rolling in hot-rolling process to prevent normal grain growth
and suppress AlN precipitation in the cooling step.
[0244] Although the cooling condition is not limited, the starting
time of the cooling step is preferably within 1.5 seconds, more
preferably 1.0 second, and most preferably 0.5 second, after the
end of finish rolling. This is because when cooling is performed
immediately after the end of rolling, a large amount of ferrite
nuclei is produced due to an increase in the degree of over cooling
with accumulated strain to promote ferrite transformation and
suppress the diffusion of dissolved N in the .gamma. phase into the
ferrite grains, thereby increasing the amount of dissolved N
present in the ferrite grain boundaries.
[0245] The cooling rate is preferably 10.degree. C./s or more in
order to ensure dissolved N. Particularly, when the finishing
temperature of hot-rolling is the Ar.sub.3 transformation point or
more, the cooling rate is preferably 50.degree. C./s or more in
order to ensure dissolved N.
[0246] Then, the hot-rolled sheet is coiled. In order to coarsen a
carbide, the coiling temperature is advantageously as high as
possible. However, with a coiling temperature of over 800.degree.
C., the scale formed on the surface of the hot-rolled sheet is
thickened to increase the load of the work of removing the scale,
and progress the formation of a nitride, causing a change in the
amount of dissolved N in the coil length direction. On the other
hand, with a coiling temperature of less than 400.degree. C., the
coiling work becomes difficult. Therefore, the coiling temperature
of the hot-rolled sheet must be in the range of 400 to 800.degree.
C.
[0247] Then, the hot-rolled sheet is cold-rolled, but the rolling
reduction ratio of cold rolling must be 60 to 95%. This is because
with a rolling reduction ratio of cold rolling of less than 60%, a
high r value cannot be expected, while with a rolling reduction
ratio of over 95%, the r value is decreased.
[0248] The cold-rolled sheet subjected to cold rolling is then
recrystallized and annealed. Although the annealing method may be
either continuous annealing or batch annealing, continuous
annealing is advantageous. The continuous annealing may be
performed either in a normal continuous annealing line or in a
continuous hot-dip galvanization line.
[0249] The preferable annealing conditions include 650.degree. C.
or more for 5 seconds or more. This is because with an annealing
temperature of less than 650.degree. C., and an annealing condition
of less than 5 seconds, recrystallization is not completed to
decrease deep drawability. In order to improve deep drawability,
annealing is preferably performed in the ferrite single phase
region at 800.degree. C. or more for 5 seconds or more.
[0250] Annealing in the high-temperature .alpha.+.gamma. two-phase
region partially produces .alpha..fwdarw..gamma. transformation to
improve the r value due to the development of the {111} aggregation
structure. However, when .alpha..fwdarw..gamma. transformation
completely proceeds, the aggregation structure is made random to
decrease the r value, thereby deteriorating deep drawability.
[0251] The upper limit of the annealing temperature is preferably
900.degree. C. This is because with an annealing temperature of
over 900.degree. C., redissolution of a carbide proceeds to
excessively increase the amount of dissolved C, thereby
deteriorating the natural aging property. When
.alpha..fwdarw..gamma. transformation occurs, the aggregation
structure is made random to decrease the r value, deteriorating
deep drawability.
[0252] Furthermore, in the heating-up step in recrystallization
annealing, slow heating is performed in the temperature region from
500.degree. C. to the recrystallization temperature to sufficiently
precipitate AlN, and the like, thereby effectively decreasing the
crystal grain diameter of the steel sheet.
[0253] The temperature region in which controlled heating must be
performed is 500.degree. C., at which AlN or the like starts to
precipitate, to the recrystallization temperature The heating rate
is preferably in the range of 1 to 20.degree. C./s because with a
heating rate of over 20.degree. C./s, the sufficient amount of
precipitates cannot be obtained, while with a heating rate of less
than 1.degree. C./s, precipitates are coarsened to weaken the
effect of suppressing grain growth.
[0254] After the recrystallization annealing, temper rolling of 10%
or less may be performed for correcting the shape and controlling
surface roughness.
[0255] The cooling rate after soaking in recrystallization
annealing is preferably 10 to 50.degree. C./s. This is because with
a cooling rate of 10.degree. C./s or less, grains are grown during
cooling to coarsen the crystal grains, thereby deteriorating the
strain aging property and natural aging property. While with a
cooling rate of 50.degree. C./s or more, dissolved N does not
sufficiently diffuse into the grain boundaries, deteriorating the
natural aging property. The cooling rate is preferably 10 to
30.degree. C./s.
[0256] After the recrystallization annealing, hot-dip galvanization
and alloying by heating are performed to form a galvannealed steel
sheet as occasion demands.
[0257] The hot-dip galvanization and alloying are not limited, and
may be performed according to a conventional known method.
[0258] With a steel sheet subjected to surface treatment generally
used for steel thin sheets, such as a steel sheet (a dull-finished
steel sheet, a bright-finished steel sheet, or a steel sheet having
a specified roughness pattern formed on the surface thereof), which
is produced by temper-rolling the alloyed hot-dip galvanized steel
sheet, for improving processability and the appearance after
processing, a steel sheet having an oil film layer of antirust oil
or lubricating oil formed on the surface thereof, or the like, the
effect of the present invention can be sufficiently exhibited in
the composition range of the prevent invention.
[0259] Therefore, a cold-rolled steel sheet and a galvannealed
steel sheet can be obtained, which have excellent deep drawability
and excellent strain age hardenability, that tensile strength
increased by press forming and heat treatment.
[0260] A description will now be made of the reasons for limiting
the components of a steel sheet in the above ranges according to a
third embodiment of the present invention.
[0261] C: Less Than 0.01 Mass %
[0262] From the viewpoint of excellent deep drawability and press
moldability, C is advantageously as small as possible. Also,
redissolution of NbC proceeds in the annealing step after cold
rolling to increase the amount of dissolved C in crystal grains,
thereby easily causing deterioration in natural aging resistance.
Therefore, the C amount is preferably suppressed to less than 0.01
mass %, more preferably 0.0050 mass % or less, and most preferably
0.0030 mass % or less. In order to ensure strength and prevent
coarsening of crystal grains, the C content is preferably 0.0005%
or more.
[0263] Si: 0.005 to 1.0 Mass %
[0264] Si is a useful component for suppressing a decrease in
elongation, and improving strength. However, with a Si content of
less than 0.005 mass %, the effect of addition of Si is
insufficient, while with a Si content of over 1.0 mass %, the
deterioration of surface properties induce a decrease in
elongation. Therefore, the Si content is limited to the range of
0.005 to 1.0 mass %, and preferably the range of 0.01 to 0.75 mass
%.
[0265] Mn: 0.01 to 1.5 Mass %
[0266] Mn not only is useful as a strengthening component for
steel, but also has the function to suppress embrittlement with S
due to the formation of MnS. However, with a Mn content of less
than 0.01 mass %, the effect of addition of Mn is insufficient,
while with a Mn content of over 1.5 mass %, the deterioration of
surface properties induce a decrease in elongation. Therefore, the
Mn content is limited to the range of 0.01 to 1.5 mass %, and
preferably the range of 0.10 to 0.75 mass %.
[0267] P: 0.01 Mass % or Less
[0268] P is a solid solution strengthening element which
effectively contributes to strengthening of steel. However, with a
P content of over 0.01 mass %, deep drawability deteriorates due to
the formation of phosphide such as (FeNb).sub.xP or the like.
Therefore, P is limited to 0.10 mass % or less.
[0269] S: 0.01 Mass % or Less
[0270] With a high S content, the amount of inclusions is increased
to deteriorate ductility. Therefore, contamination with S is
preferably prevented as much as possible, but an S content up to
0.01 mass % is allowable.
[0271] Al: 0.005 to 0.030 Mass %
[0272] Al is added as a element for deoxidization for improving the
yield of the elements forming carbonitride. However, with an Al
content of less than 0.005 mass %, the effect is insufficient,
while with an Al content of over 0.030 mass %, the amount of N to
be added to steel is increased to easily cause slab defects during
steel making. Therefore, Al is contained in the range of 0.005 to
0.030 mass %.
[0273] N: 0.005 to 0.040 Mass %
[0274] In the present invention, N is an important element which
plays the role of imparting strain age hardenability to a steel
sheet. However, with an N content of less than 0.005 mass %, a
sufficient strain age hardenability cannot be obtained, while with
an N content of as high as over 0.040 mass %, press moldability
deteriorates. Therefore, N is contained in the range of 0.005 to
0.040 mass %, and preferably in the range of 0.008 to 0.015 mass
%.
[0275] B: 0.0001 to 0.003 Mass %
[0276] B is added in a combination with Nb to exhibit the function
to effectively make fine the micro structure of the hot-rolled
steel and the cold-rolled steel, annealed for recrystallization,
and to improve cold-work embrittlement resistance. However, with a
B content of less than 0.0001 mass %, the sufficient effect of
making fine the structures cannot be obtained, while with a B
content of over 0.003 mass the amount of BN precipitate is
increased, and dissolution in the slab heating step is hindered.
Therefore, B is contained in the range of 0.0001 to 0.003 mass %,
preferably in the range of 0.0001 to 0.0015 mass %, and more
preferably in the range of 0.0007 to 0.0012 mass %.
[0277] Nb: 0.005 to 0.050%, Ti: 0.005 to 0.070%, V: 0.005 to
0.10%
[0278] Nb, Ti and V are added in a combination with B to contribute
to refinement of the the micro structure of the hot-rolled steel
and the cold-rolled steel, annealed for recrystallization, and have
the function to precipitate dissolved C as NbC, Tic and VC,
respectively. Therefore, these elements are added together with B
according to demand, but less than 0.005% each of the elements does
not sufficiently exhibit the function. On the other hand, over
0.050% of Nb, over 0.070% of Ti and over 0.10% of V cause
deterioration in ductility. Therefore, Nb, Ti and V are added in
the ranges of 0.005 to 0.050%, 0.005 to 0.070%, and 0.005 to 0.10%,
respectively.
[0279] Furthermore, as described above, Nb has the function to fix
dissolved C as NbC, and forms a nitride NbN. Similarly, Al and B
form AlN and NB, respectively. Therefore, in order to ensure the
sufficient amount of dissolved N and sufficiently decrease the
amount of dissolved C, it is important to satisfy the following
relations (1) and (2):
N %.gtoreq.0.0015+14/93.multidot.Nb %+14/27.multidot.Al
%+14/11.multidot.B % (1)
C %.ltoreq.0.5.multidot.(12/93).multidot.Nb % (2)
[0280] N/Al or N/(Al+Nb+Ti+V+B): 0.30 or More
[0281] Al forms AlN to decrease the amount of dissolved N. In order
to ensure an appropriate amount of dissolved N, N/Al must be 0.30
or more. When Al is added in a combination with Nb, Ti, V or B,
these elements also respectively form NbN, TiN, VN and BN to
decrease the amount of dissolved N. Therefore, in order to ensure
an appropriate amount of dissolved N, (Al+Nb+Ti+V+B) must be 0.30
or more.
[0282] Dissolved N: 0.0010% or More
[0283] In order to increase the strain age hardenability of the
steel sheet, the content of dissolved N must be 0.0010% or
more.
[0284] The amount of dissolved N is determined by subtracting the
amount of precipitated N from the total N amount of steel. As a
result of research of comparison between various analysis methods,
the inventors found that an electrolytic extraction method using a
constant-potential electrolytic method is effective as the method
of analyzing the amount of precipitated N. As a method of
dissolving ferrite used for extraction analysis, an acid digestion
method, a halogen method, or an electrolysis method can be used. Of
these methods, the electrolysis method can stably dissolve only
ferrite without decomposing very unstable precipitates such as a
carbide, a nitride, etc. As the electrolyte, an acetyl-acetone
system is used for electrolysis at a constant potential. In the
present invention, the results of measurement of the amount of
precipitated N by constant-potential electrolysis showed best
correspondence with the actual strength of parts.
[0285] Therefore, in the present invention, the residue after
extraction by constant-potential electrolysis is chemically
analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated
N.
[0286] In order to obtain higher BH and .DELTA.TS, the amount of
dissolved N is preferably 0.0015% or more, more preferably 0.0020%
or more, and most preferably 0.0030% or more.
[0287] The cold-rolled steel sheet of the present invention is a
cold-rolled deep drawing steel sheet having excellent strain age
hardenability and the above-described composition, wherein
TS.times.r value.gtoreq.750 MPa.
[0288] A steel sheet having a Ts x r value of less than 750 MPa
cannot be widely applied to members comprising structural member
components. In order to extend the application range, the
TS.times.r value is preferably 850 MPa or more.
[0289] Conventional coating and baking conditions include
170.degree. C. for 20 min as standards. When a strain of 5% is
applied to the steel sheet of the present invention, which contains
a large amount of dissolved N, hardening can be achieved even by
slow (low temperature) processing. In other words, the range of
aging conditions can be widened. In order to attain a sufficient
amount of hardening, generally, retention at a higher temperature
for a longer time is advantageous as long as softening does not
occurs by over aging.
[0290] Specifically, in the steel sheet of the present invention,
the lower limit of the heating temperature at which hardening
significantly takes place after pre-deformation is about
100.degree. C. On the other hand, with the heating temperature of
over 300.degree. C., hardening peaks, thereby causing the tendency
to soften and significantly causing thermal strain and temper
color. With the retention time of about 30 seconds or more,
hardening can be sufficiently achieved at a heating temperature of
about 200.degree. C. In order to obtain more stable hardening, the
retention time is preferably 60 seconds or more. However, retention
for over 20 mines is practically disadvantageous because further
hardening cannot be expected, and the production efficiency
significantly deteriorates.
[0291] Therefore, in the present invention, the conventional
coating and baking conditions, i.e., the heating temperature of
170.degree. C. and the retention time of 20 minutes, are set as the
aging conditions. With the steel sheet of the present invention,
hardening can be stably achieved even under the aging conditions of
a low heating temperature and a short retention time, which fail to
achieve sufficient hardening in a conventional bake-hardening steel
sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and
other methods such as induction heating, heating with a
nonoxidation flame, a laser, plasma, or the like, etc. can be
preferably used. Alternatively, only a portion in which strength is
desired to be increased may be selectively heated.
[0292] The strength of an automobile part must be sufficient to
resist an external complicated stress load, and thus not only
strength in a low strain region but also strength in a high strain
region are important for a raw material steel sheet. In
consideration of this point, in the steel sheet of the present
invention used as a raw material for automobile parts, BH is 80 MPa
or more, and .DELTA.TS is 40 MPa or more. More preferably, BH is
100 MPa or more, and .DELTA.TS is 50 MPa or more. In order to
further increase BH and .DELTA.TS, the heating temperature in aging
may be set to a higher temperature, and/or the retention time may
be set to a longer time.
[0293] The steel sheet of the present invention has the advantage
that even when the steel sheet not molded is allowed to stand at
room temperature for about one year, natural aging deterioration
does not occur, unlike a conventional steel sheet.
[0294] In the present invention, the cold-rolled steel sheet may be
coated by hot-dip galvanization or alloying hot-dip galvanization
without any problem, and TS, BH and .DELTA.TS are equivalent to
those before plating. Besides hot-dip galvanization and alloying
hot-dip galvanization, electro-galvanization, electro-tinning,
electric chromium plating, electro-nickeling, and the like may be
preferably used.
[0295] The production conditions according to a fourth embodiment
of the present invention will be described.
[0296] First, steel having the composition, by mass %, comprising
less than 0.01% of C, 0.0050 to 0.04% of N, 0.005 to 0.03% of Al,
0.005 to 1.0% of Si, 0.01 to 1.5% of Mn, 0.1% or less of P, and
0.01% or less of S, or further comprising 0.0001 to 0.003% of B,
and at least one of 0.005 to 0.050% of Nb, 0.005 to 0.070% of Ti,
and 0.005 to 0.10% of V, wherein N/(Al+Nb+Ti+V+B) is 0.30 or more,
is melted by a conventional melting method such as a converter or
the like, and then solidified by an ingot making method or a
continuous casting method to form a steel raw material.
[0297] The steel raw material is heated and soaked, and then
hot-rolled to form a hot-rolled sheet. With an excessively low
heating temperature (SRT), the effect of improving processability
is saturated, and the rolling load in hot rolling is increased to
cause a trouble in rolling, and insufficient homogeneity of
dissolved N. Therefore, SRT is preferably 950.degree. C. or more.
In order to improve deep drawability, it is advantageous to fix
dissolved C and precipitate it as a carbide. Therefore, SRT is
preferably 1300.degree. C. or less. In order to further improve
processability, SRT is preferably 1150.degree. C. or less.
[0298] When the total rolling reduction ratio of hot rolling
starting from rough rolling to finish rolling is less than 80%, the
crystal grains of the hot-rolled sheet are not sufficiently made
fine. Therefore, the total rolling reduction ratio is preferably
80% or more.
[0299] With the rough rolling temperature of over 1000.degree. C.,
.gamma..fwdarw..alpha. transformed grains are coarsened to decrease
the r value, while with the rough rolling temperature of less than
Ar.sub.3 transformation point, a grains are recrystallized and
coarsened or grown to decrease the r value. Therefore, rough
rolling is preferably performed in the temperature region of the
Ar.sub.3 transformation point to 1000.degree. C.
[0300] On the other hand, when finish rolling is completed in the
temperature region of over the Ar.sub.3 transformation point, the
aggregation structure is made random by .gamma..fwdarw..alpha.
transformation to fail to obtain excellent deep drawability. When
finish rolling is completed in the temperature region of less than
the Ar.sub.3 transformation point, a further improvement in deep
drawability cannot be expected, but the rolling load is increased.
Therefore, finish rolling is preferably performed in the
temperature region of 600.degree. C. to the Ar.sub.3 transformation
point.
[0301] If lubrication rolling is not performed in finish rolling,
an additional shear force acts on the surface of the steel sheet
due to frictional force between a roll and the steel sheet to
preferentially form {110} orientation undesirable for deep drawing
in the surface of the steel sheet, deteriorating deep drawability.
Therefore, finish rolling is preferably performed under
lubrication.
[0302] Then, the hot-rolled sheet is coiled. The processed material
after the coiling step is referred to as a "coil". The higher
coiling temperature (CT) of the hot-rolled sheet is advantageous
for coarsening of the carbide. However, with the coiling
temperature of over 800.degree. C., the scale formed on the surface
of the hot-rolled sheet is thickened to increase the load of a
scale removing work, and progress the formation of a nitride,
thereby causing a variation in amount of dissolved N in the
longitudinal direction of the coil. On the other hand, with the
coiling temperature of less than 400.degree. C., the coiling work
is difficult. Therefore, CT is preferably 400 to 800.degree. C.
[0303] Then, the thus-obtained hot-rolled sheet is recrystallized
and annealed by continuous annealing or batch annealing. The
annealing (hot-rolled sheet annealing) is carried out for
recrystallizing the rolled aggregation structure formed by hot
rolling in the a-phase region in finish rolling to obtain a
recrystallized aggregation structure.
[0304] Then, the hot-rolled sheet is cold-rolled to form a
cold-rolled sheet. When the rolling reduction ratio of cold rolling
is less than 60%, a high r value cannot be expected. On the other
hand, while a rolling reduction ratio of over 95%, the r value is
decreased. Therefore, the rolling reduction ratio is preferably 60
to 95%.
[0305] Next, the cold-rolled sheet is recrystallized and annealed.
The annealing is preferably carried out in either a continuous
annealing line or a continuous hot-dip galvanization line. The
preferable annealing conditions include the annealing temperature
of 650.degree. C. or more and the retention time of 5 seconds or
more. When either of the annealing temperature of 650.degree. C. or
more and the retention time of 5 seconds or more is not satisfied,
recrystallization is not completed to deteriorate deep drawability.
In order to obtain excellent deep drawability, the annealing
temperature of 800.degree. C. or more and the retention time of 5
seconds or more are preferred. However, with the annealing
temperature of over 900.degree. C., redissolution of the carbide
proceeds to excessively increase the amount of dissolved C, thereby
deteriorating the natural aging property (decreasing elongation by
natural aging). Furthermore, when .gamma..fwdarw..alpha.
transformation occurs, the aggregation structure becomes random to
decrease the r value, deteriorating deep drawability. Therefore,
the annealing temperature is preferably 900.degree. C. or less.
[0306] The cold-rolled annealed sheet obtained by recrystallizing
and annealing the cold-rolled steel sheet is further coated by
hot-dip galvanization or alloyed. In this case, in plating, the
cooling rate during the time between the completion of
recrystallization annealing and the start of plating is 5.degree.
C./s or more, and the sheet temperature in hot-dip galvanization is
preferably 400 to 600.degree. C. In alloying, the processing
temperature is preferably 400 to 600.degree. C., and the processing
time is preferably 5 to 40 seconds.
[0307] The cold-rolled steel sheet after recrystallization
annealing or the hot-dip galvanized steel sheet may be
temper-rolled for correcting the shape and controlling surface
roughness. The reduction ratio of temper rolling is preferably 10%
or less. This is because with a rolling reduction ratio of over
10%, the r value is decreased.
[0308] Description will now be made of the reasons for limiting the
composition of a high-tensile-strength cold-rolled steel sheet
according to a fifth embodiment of the present invention.
[0309] C: 0.0015 to 0.025%
[0310] In the present invention, in order to control the structure
to a homogeneous fine structure, and ensure a sufficient amount of
an acicular ferrite phase, the C content must be 0.0015% or more.
With a C content of over 0.025%, the ratio of the carbide in the
steel sheet is excessively increased to significantly deteriorate
ductility, the r value and moldability. Therefore, the C content is
limited in the range of 0.0015 to 0.025%. From the viewpoint of
improvement in moldability, the C content is preferably 0.020% or
less, more preferably 0.010% or less. Particularly, from the
viewpoint of stabilization of the BH amount and material
properties, the C content preferably exceeds (12/93) Nb (%)
(wherein Nb represents the Nb content (%)).
[0311] Si: 1.0% or Less
[0312] Si is a useful component capable of increasing the strength
of the steel sheet without significantly deteriorating ductility of
steel. Particularly, when high strength is required, the Si content
is preferably 0.10% or more. On the other hand, Si is an element
which greatly changes the transformation point during hot rolling
to cause difficulties in ensuring quality and the shape, or
adversely affects surface properties, chemical conversion
properties, and the like, particularly the beauty of the surface of
the steel sheet, and adversely affects plating properties. In the
present invention, therefore, the Si content is limited to 1.0% or
less. However, the above-described adverse effects can be kept down
as long as Si is 1.0% or less. Particularly, in applications
required for the steel sheet to have surface beauty, Si is
preferably 0.5% or less.
[0313] Mn: 2.0% or Less
[0314] Mn is an element effective to prevent hot cracking with S,
and Mn is preferably added according to the amount of S contained.
Mn also has the great effect of making fine crystal grains, and is
preferably added for improving material properties. In order to
stably fix S, the Mn content is preferably 0.1% or more. Mn is also
an element for increasing the strength of the steel sheet, and is
preferably added in an amount of 0.5% or more when higher strength
is required. The Mn content is more preferably 0.8% or more.
[0315] With the Mn content increased to this level, there is the
advantage that variations in the mechanical properties of the steel
sheet with respect to variations in the hot-rolling conditions,
particularly strain age hardenability, are significantly improved.
However, with the excessively high Mn content of over 2.0%, the
deformation resistance at elevated temperatures tends to increase
to deteriorate weldability and weld moldability. The detailed
mechanism of this is not known. Furthermore, the formation of
ferrite is significantly suppressed, and the r value is
significantly decreased. Therefore, the Mn content is limited to
2.0% or less. In applications required to have good corrosion
resistance and moldability, the Mn content is preferably 1.5% or
less.
[0316] P: 0.1% or Less
[0317] P is a useful element as a solid solution strengthening
element for steel, and is preferably added in an amount of 0.002%
or more from the viewpoint of an increase in strength.
Particularly, when high strength is required, the P content is
preferably 0.02% or more. On the other hand, when P is excessively
added, steel is embrittled, and stretch-flanging properties of the
steel sheet deteriorate. Also, P is liable to strongly segregate in
steel, thereby causing embrittlement of a weld. Therefore, P is
limited to 0.1% or less. In applications in which elongated flange
processability and weld toughness are considered as important, P is
preferably 0.08% or less, more preferably 0.06% or less.
[0318] S: 0.02% or Less
[0319] S is present as an inclusion in the steel sheet, decreases
ductility of the steel sheet, and causes deterioration in corrosion
resistance. Therefore, the Si content is as low as possible, and in
the present invention, the Si content is limited to 0.02% or less.
Particularly, in applications required to have good processability,
S is preferably 0.015% or less. Particularly, in applications
required to have excellent stretch-flanging properties, S is
preferably 0.010% or less. Although the detailed mechanism is not
known, in order to stably maintain the strain age hardenability of
the steel sheet in a high level, it is effective to decrease the S
content to 0.008% or less.
[0320] Al: 0.02% or Less
[0321] Al is an element functioning as a element for deoxidation
for improving cleanliness of steel, and making fine the structure
of the steel sheet. In the present invention, the Al content is
preferably 0.001% or more. In the present invention, dissolved N is
used as a strengthening element, but aluminum killed steel
containing Al in a suitable range has mechanical properties
superior to those of conventional rimmed steel not containing Al.
On the other hand, with an excessively high Al content, the surface
properties of the steel sheet deteriorate, and the amount of
dissolved N is significantly decreased to cause difficulties in
obtaining a large amount of strain age hardening, which is the main
object of the present invention. Therefore, in the present
invention, Al is limited to 0.02% or less. From the viewpoint of
stability of material quality, Al is more preferably 0.001 to
0.015%. Although a decrease in the Al content possibly causes
coarsening of crystal grains, in the present invention, the amounts
of other alloy elements are appropriately determined, and the
annealing conditions are set in appropriately ranges, thereby
effectively preventing coarsening.
[0322] N: 0.0050 to 0.0250 Mass %
[0323] N is an element for increasing the strength of the steel
sheet by solid solution strengthening and strain age hardening, and
in the present invention, N is the most important element. In the
present invention, an appropriate amount of N is contained, the Al
content is controlled to the appropriate value, and production
conditions such as the hot-rolling conditions, and the annealing
conditions are controlled to ensure necessary and sufficient
dissolved N in a cold-rolled product or a coated product. This
exhibits the sufficient effect of increasing strength (yield stress
and tensile strength) by solid solution strengthening and strain
age hardening, to stably obtain the target values of the mechanical
properties of the steel sheet of the present invention, such as a
tensile strength of 340 MPa or more, an amount (BH amount) of
bake-hardening of 80 MPa or more, and an increase is tensile
strength .DELTA.TS of 40 MPa or more after strain aging. Since N
also has the function to decrease the transformation point, N is
effective for rolling of a thin material for which rolling at a
temperature greatly over the transformation point is
undesirable.
[0324] With a N content of less than 0.0050%, the effect of
increasing strength is less stably exhibited, while with a N
content of over 0.0250%, the rate of occurrence of internal defects
in the steel sheet is increased, and slab cracking frequently
occurs during continuous casting. Therefore, N is limited to the
range of 0.0050 to 0.0250%. From the viewpoint of improvement in
stability of material properties and yield over the entire
production process, N is preferably in the range of 0.0070 to
0.0200%, and more preferably in the range of 0.0100 to 0.0170%.
With the N amount in the range of the present invention, there is
no adverse effect on weldability, and the like.
[0325] Dissolved N: 0.0010% or More
[0326] In order to ensure sufficient strength of a cold-rolled
product, and effectively exhibit strain age hardening with N, it is
necessary that the content of dissolved N (solid solution N) in the
steel sheet is at least 0.0010% or more.
[0327] The amount of dissolved N is determined by subtracting the
amount of precipitated N from the total N amount of steel. As a
result of comparison research of various methods, the inventors
found that electrolytic extraction analysis using
constant-potential electrolysis is effective as the method of
analyzing the amount of precipitated N. As the method of dissolving
ferrite used for extraction analysis, an acid digestion method, a
halogen method, or an electrolysis method can be used. Of these
methods, the electrolysis method can stably dissolve only ferrite
without decomposing very unstable precipitates such as a carbide, a
nitride, etc. As the electrolyte, an acetyl-acetone system is used
for electrolyzing at a constant potential. In the present
invention, the results of measurement of the amount of precipitated
N by constant-potential electrolysis showed best correspondence
with changes in actual material properties.
[0328] Therefore, in the present invention, the residue after
extraction by constant-potential electrolysis is chemically
analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated
N.
[0329] In order to obtain higher BH and .DELTA.TS, the amount of
dissolved N is preferably 0.0020% or more. In order to obtain
further high values, the amount of dissolved N is preferably
0.0030% or more. Although the upper limit of the amount of
dissolved N is not limited, the mechanical properties less
deteriorate even when the all amount of N remains.
[0330] N/Al (the Content Ratio of N to Al): 0.3 or more
[0331] In order to cause 0.0010% or more of dissolved N to stably
remain in a product state, the amount of Al which is an element for
strongly binding to N, must be limited. As a result of research of
steel sheets in which the combination of the N content (0.0050 to
0.0250%) and the Al content (0.02% or less) were widely changed in
the composition range of the present invention, it was found that
with N/Al of 0.3 or more, the amount of dissolved N of a
cold-rolled product or coated product can be stably set to 0.0010%
or more. Therefore, N/Al is limited to 0.3 or more. In order to
stably increase the strain age hardenability, N/Al is preferably
0.6 or more, and more preferably 0.8 or more.
[0332] Nb: 0.002 to 0.050%
[0333] Nb effectively functions to form an acicular ferrite phase
in combination with B. In the present invention, the Nb content
must be 0.002% or more. On the other hand, a Nb content of over
0.050%, the effect is saturated, and deformation resistance at
elevated temperatures is significantly increased to cause
difficulties in hot rolling. Therefore, Nb is limited in the range
of 0.002 to 0.050%, and preferably 0.005 to 0.040%.
[0334] B: 0.0001 to 0.0050%
[0335] B effectively functions to form an acicular ferrite phase in
combination with Nb. In the present invention, the B content must
be 0.001% or more. On the other hand, a B content of over 0.0050%,
the amount of dissolved N contributing strain age hardenability is
decreased. Therefore, B is limited in the range of 0.0001 to
0.0050%, and preferably 0.0003 to 0.0030%.
[0336] In the present invention, the above composition preferably
further contains at least one of the following groups a to c:
[0337] Group a: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0338] Group b: one or both of Ti and V in a total of 0.1% or less;
and
[0339] Group c: one or both of Ca and REM in a total of 0.0010 to
0.010%.
[0340] Element of group a: Cu, Ni, Cr and Mo are all contribute to
an increase in strength of the steel sheet, and can be contained
singly or in a combination according to demand. The effect is
recognized by containing 0.01% or more each of Cu, Ni, Cr and Mo.
However, with an excessively high content, deformation resistance
at elevated temperatures in hot rolling is increased, or chemical
conversion properties and surface treatment properties in a wide
sense deteriorate, and a welded portion is hardened to deteriorate
weld moldability. Therefore, Cu, Ni, Cr, and Mo are preferably
contained singly at 1.0% or less, 1.0% or less, 0.5% or less, and
0.2% or less, respectively, and preferably contained in a
combination at a total of 1.0% or less.
[0341] Element of group b: Both Ti and V are elements contributing
refinement and homegenization of crystal grains, and may be added
singly or in a combination according to demand. The effect can be
recognized by containing 0.005% or more each of Ti and V. However,
with an excessively high content, deformation resistance at
elevated temperatures in hot rolling is increased, or chemical
conversion properties and surface treatment properties in a wide
sense deteriorate. Furthermore, there is the adverse effect of
decreasing the amount of dissolved N. Therefore, Ti and V are
preferably contained singly at 1.0% or less and 1.0% or less,
respectively, and preferably contained in a combination at a total
of 0.1% or less.
[0342] Elements of group c: Both Ca and REM are elements useful for
controlling the form of inclusions. Particularly, when the stretch
flanging property is required, these elements are preferably added
singly or in a combination. When the total of the elements of group
d is less than 0.0010%, the effect of controlling the form of
inclusions is insufficient, while when the total exceeds 0.010%,
surface defects significantly occur. Therefore, the total of the
elements of group d is preferably limited to the range of 0.0010 to
0.010%. This permits improvement in the stretch flanging property
without causing surface defects.
[0343] The structure of the steel sheet of the present invention is
described below.
[0344] The steel sheet of the present invention has the structure
composed of an acicular ferrite phase at an area ratio of 5% or
more and a ferrite phase having an average crystal grain diameter
of 20 .mu.m or less.
[0345] The Area Ratio of the Acicular Ferrite Phase: 5% or More
[0346] The cold-rolled steel sheet of the present invention
contains the acicular ferrite phase at an area ratio of 5% or more.
The presence of the acicular ferrite phase at 5% or more permits
the achievement of good ductility and a larger amount of strain age
hardening. Although the detailed mechanism is not known, it is
thought that strain is effectively accumulated in the steel sheet
during pre-strain processing before aging by the presence of the
acicular ferrite phase. Furthermore, the presence of the acicular
ferrite phase improves natural aging deterioration at room
temperature to be effective to obtain natural non-aging properties.
In order to obtain a good balance of strength and ductility, and
higher strength, the area ratio of the acicular ferrite phase is
preferably 10% or more. The presence of a large amount of acicular
ferrite phase of over 20% has the problem of deteriorating the r
value. Therefore, the area ratio of the acicular ferrite phase is
5% or more, and preferably 10% to 20%.
[0347] In the present invention, "the acicular ferrite phase" is a
low-temperature transformation phase peculiar to ultra low carbon
steel having the composition of the present invention, in which no
carbide is contained therein. This phase can be clearly
discriminated from normal polygonal ferrite by observation on an
optical microscope, and is harder than the polygonal ferrite phase
because of the high internal dislocation density.
[0348] According to observation on an optical microscope, the
acicular ferrite phase has a distribution of any one of (1) crystal
grains having irregularly angular boundaries, (2) crystal grains
present along crystal grains of precipitates or the like, and (3)
crystal grains or crystal grain groups (many sub-boundaries are
observed in relatively second phase grains) having a scratch-like
pattern, singly or in a combination. This acicular ferrite can be
clearly distinguished from general polygonal ferrite. The color
tone of the corroded insides of the grains is different from
martensite and bainite, and substantially the same as ordinary
polygonal ferrite, and thus acicular ferrite can be clearly
distinguished from martensite and bainite. According to observation
on a transmission electron microscope, the acicular ferrite phase
has a very high density of dislocation in the vicinities of the
crystal grains and/or in the grains, and particularly the above
form (3) comprises a layer having a very high dislocation density
and a layer having a relatively low dislocation density.
[0349] The cold-rolled steel sheet of the present invention is
directed to use as an automobile steel sheet required to have high
moldability, and comprises a ferrite phase other than the acicular
ferrite phase in order to ensure ductility. When the area ratio of
the ferrite phase is less than 80%, it is difficult to ensure
ductility and a high r value which are necessary for the automobile
steel sheet required to have processability. When further improved
ductility is required, the area ratio of the ferrite phase is 80%
or more, and preferably 85% or more. In the present invention,
ferrite means so-called polygonal ferrite in which no strain
remains.
[0350] Average Crystal Grain Diameter of Ferrite Phase: 20 .mu.m or
Less
[0351] In the present invention, the value used as the average
crystal grain diameter is a higher one of the value calculated from
a photograph of a sectional structure by a quadrature method
defined by ASTM, and the nominal value determined by an intercept
method defined by ASTM (refer to, for example, Umemoto et al.: Heat
Treatment, 24 (1984), p334).
[0352] The cold-rolled steel sheet of the present invention
maintains a predetermined amount of dissolved N in the product
step. However, as a result of experiment and research conducted by
the inventors, it was found that variations in strain age
hardenability occur in steel sheets containing the same amount of
dissolved N, and one of the main causes of the variations is a
crystal grain diameter. In the structure of the present invention,
in order to stably obtain a high BH amount and .DELTA.TS, the
average crystal grain diameter is at least 20 .mu.m or less, and
preferably 15 .mu.m or less. Although the detailed mechanism is not
known, this is supposed to be related to the segregation and
precipitation of alloy elements in the crystal grain boundaries,
and the influences of processing and heat history on the
segregation and precipitation.
[0353] Therefore, in order to achieve stability of strain age
hardenability, the average crystal grain diameter of the ferrite
phase is 20 .mu.m or less, and preferably 15 .mu.m or less.
[0354] The cold-rolled steel sheet of the present invention, which
has the above-described composition and structure, has a tensile
strength (TS) of about 340 MPa to 590 MPa, a r value of as high as
1.2 or more, an excellent strain age hardenability. The steel sheet
having TS of less than 340 MPa cannot be widely applied to members
each comprising a structural component. In order to extend the
application range, TS is preferably 400 MPa or more. With a r value
of less than 1.2, the steel sheet cannot be applied to a wide range
of press-molded products. The preferred range of the r value is 1.3
or more.
[0355] The conventional coating and baking conditions include
170.degree. C. and 20 min as standards. When a strain of 5% or more
is applied to the steel sheet of the present invention containing a
large amount of dissolved N, hardening can be achieved even by
aging at low temperature. In other words, the aging conditions can
be selected from a wide range. Generally, in order to earn the
amount of hardening, it is advantageous to hold the steel sheet at
a higher temperature for a longer time as long as softening does
not occur due to over aging.
[0356] Specifically, in the steel sheet of the present invention,
the lower limit of the heating temperature at which hardening
significantly takes place after pre-deformation is about
100.degree. C. On the other hand, with the heating temperature of
over 300.degree. C., hardening peaks, thereby causing the tendency
to soften and significantly causing thermal strain and temper
color. With the retention time of about 30 seconds or more,
hardening can be sufficiently achieved at a heating temperature of
about 200.degree. C. In order to obtain more stable hardening, the
retention time is preferably 60 seconds or more. However, retention
for over 20 minutes is practically disadvantageous because further
hardening cannot be expected, and the production efficiency
significantly deteriorates.
[0357] Therefore, in the present invention, the conventional
coating and baking conditions, i.e., the heating temperature of
170.degree. C. and the retention time of 20 minutes, are set as the
aging conditions. With the steel sheet of the present invention,
hardening can be stably achieved even under the aging conditions of
a low heating temperature and a short retention time, which fail to
achieve sufficient hardening in a conventional bake-hardening steel
sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and
other methods such as induction heating, heating with a
nonoxidation flame, a laser, plasma, or the like, etc. can be
preferably used.
[0358] The strength of an automobile part must be sufficient to
resist an external complicated stress load, and thus not only
strength in a low strain region but also strength in a high strain
region are important for a raw material steel sheet. In
consideration of this point, in the steel sheet of the present
invention used as a raw material for automobile parts, BH is 80 MPa
or more (corresponding to strength in a relatively low strain
region), and .DELTA.TS is 40 MPa or more (corresponding to strength
in a relatively high strain region). More preferably, BH is 100 MPa
or more, and .DELTA.TS is 50 MPa or more. In order to further
increase BH and .DELTA.TS, the heating temperature in aging may be
set to a higher temperature, and/or the retention time may be set
to a longer time.
[0359] The effect of the present invention is exhibited by a
product having a relatively large thickness. However, with a
product having a thickness of over 3.2 mm, a sufficient cooling
rate necessary for the cold-rolled sheet annealing step cannot be
ensured to cause strain aging at the time of continuous annealing,
thereby failing to obtain the target strain age hardenability as a
product. Therefore, the steel sheet of the present invention
preferably has a thickness of 3.2 mm or less.
[0360] In the present invention, the surface of the cold-rolled
steel sheet may be coated by hot-dip galvanization or alloying
hot-dip galvanization without any problem. These coated steel
sheets also exhibit TS, BH and .DELTA.TS which are equivalent to
those before plating. As the plating type, any one of
electro-galvanization, hot-dip galvanization, alloying hot-dip
galvanization, electro-tinning, electric chromium plating,
electro-nickeling, and the like may be preferably used.
[0361] The method of producing a steel sheet according to a sixth
embodiment of the present invention is described.
[0362] The steel sheet of the present invention is basically
produced by performing the hot rolling step in which a steel slab
having the above-described composition is heated, and then roughly
rolled to form a sheet bar, and the sheet bar is finish-rolled and
cooled to form a coiled hot-rolled sheet, the cold rolling step in
which the hot-rolled sheet is pickled and cold-rolled to form a
cold-rolled sheet, and the cold-rolled sheet annealing step in
which the cold-rolled sheet is continuously annealed.
[0363] Although the slab used in the production method of the
present invention is preferably produced by a continuous casting
method in order to prevent macro-segregation of components, an
ingot making method or a thin slab casing method may be used.
Alternatively, a conventional method comprising cooling the
produced slab to room temperature and then again heating the slab,
or an energy-saving process of direct rolling, in which a hot slab
is charged into a heating furnace without cooling and then rolled,
or the slab is rolled directly immediately after being slightly
kept warm, may be used without a problem. Particularly, direct
rolling is a useful technique for effectively ensuring dissolved
N.
[0364] Description will be now be made of the reasons for limiting
the conditions of the hot rolling step.
[0365] Slab Heating Temperature: 1000.degree. C. or More
[0366] The slab heating temperature is preferably 1000.degree. C.
or more in order to ensure a necessary and sufficient amount of
dissolved N in an initial state and satisfy the target amount of
dissolved N in a product. Since a loss is increased by an increase
in the oxide weight, the heating temperature is preferably
1280.degree. C. or less.
[0367] The slab heated under the above condition is roughly rolled
to form a sheet bar. The condition of the rough rolling is not
defined, and rough rolling may be performed according to a normal
method. However, in order to ensure the amount of dissolved N,
rough rolling is preferably performed within a as short time as
possible. Then, the sheet bar is finish-rolled to form a hot-rolled
sheet.
[0368] In the present invention, the adjacent sheet bars are
preferably bonded together during the time between rough rolling
and finish rolling, and then continuously rolled. As the bonding
means, a pressure welding method, a laser welding method, an
electron beam welding method, or the like is preferably used.
[0369] By continuous rolling, so-called non-stationary portions at
the front end and the rear end of a coil (processed material) are
removed to permit hot rolling over the entire length and the entire
width of the coil (processed material) under stable conditions.
This is very effective to improve the sectional shape and
dimensions of not only the hot-rolled steel sheet but also the
cold-rolled steel sheet. When the steel sheet is cooled on a hot
run table after rolling, the shape of the steel sheet can be
sufficiently maintained because tension can be always applied.
[0370] By continuous rolling, the ends of the coil can be stably
passed, and it is thus possible to use lubricating rolling, which
cannot be easily applied to ordinary single rolling for each sheet
bar because of the problem of continuous rolling processes and
biting property. Therefore, the rolling load can be decreased, and
at the same time, the surface pressure of the roll can be
decreased, thereby increasing the life time of the roll.
[0371] In the present invention, at the entrance of a finisher
between rough rolling and finish rolling, the temperature
distributions of the sheet bar in the width direction and the long
direction thereof are preferably made uniform by using any one or
both of a sheet bar edge heater for heating the ends of the sheet
bar in the width direction and a sheet bar heater for heating the
ends of the sheet bar in the long direction. This can further
decrease the variations in material properties of the steel sheet.
The sheet bar edge heater and the sheet bar heater are preferably
of an induction heating type.
[0372] As the procedure of use of the heaters, preferably, the
difference in temperature in the width direction is first corrected
by the sheet bar edge heater. Although the heating amount depends
upon the steel composition, the temperature distribution in the
width direction at the finisher entrance is preferably set in the
range of about 20.degree. C. or less. Next, the difference in
temperature in the long direction is corrected by the sheet bar
heater. At this time, the heating amount is preferably set so that
the temperatures at the ends in the long direction are about
20.degree. C. higher than the temperature at the center.
[0373] Finisher Delivery Temperature: 800.degree. C. or more
[0374] In order to obtain a homogeneous fine hot-rolled base sheet
structure, the finisher deliver temperature FDT is 800.degree. C.
or more. With a FDT of lower than 800.degree. C., the structure of
the steel sheet becomes inhomogeneous, and the processed structure
partially remains to leave heterogeneity of the structure after the
cold-rolled sheet annealing step. Therefore, the danger of
producing various troubles in press forming is increased. When a
high coiling temperature is used for avoiding the processed micro
structure from remaining, coarse crystal grains are produced to
cause the same troubles as described above. With the high coiling
temperature, the amount of dissolved N is significantly decreased
to cause difficulties in obtaining a target tensile strength of 340
MPa or more. Therefore, the finisher deliver temperature FDT is
800.degree. C. or more. In order to further improve the mechanical
properties, the FDT is preferably 820.degree. C. or more. In order
to improve the r value, the FDT is preferably the Ac.sub.3
transformation point or more. Although the upper limit of FDT is
not limited, a scale scar significantly occurs at excessively high
FDT. The FDT is preferably up to about 1000.degree. C.
[0375] Coiling Temperature: 800.degree. C. or Less
[0376] The strength of the steel sheet is liable to increase as the
coiling temperature CT decreases. In order to ensure the target
tensile strength TS of 340 MPa or more, the CT is preferably
800.degree. C. or less. With a CT of less than 200.degree. C., the
shape of the steel sheet is readily disturbed to increase the
danger of causing troubles in a practical operation, thereby
deteriorating homogeneity of material properties. Therefore, the CT
is preferably 200.degree. C. or more. When the homogeneity of the
material properties is required, the CT is preferably 300.degree.
C. or more, and more preferably 350.degree. C. or more.
[0377] In the present invention, in finish rolling, lubricating
rolling may be performed for decreasing the hot rolling load. The
lubricating rolling has the effect of further making homogeneous
the shape and material properties of the hot-rolled sheet. During
the lubricating rolling, the frictional coefficient is preferably
in the range of 0.25 to 0.10. By combining lubricating rolling and
continuous rolling, the operation of hot rolling is further
stabilized.
[0378] The hot-rolled sheet subjected to the above-described hot
rolling step is then pickled and cold-rolled in the cold rolling
step to form a cold-rolled sheet.
[0379] The pickling conditions may be the same as conventional
known conditions, and are not limited. When the scale of the
hot-rolled sheet is extremely small, cold rolling may be
immediately after hot rolling without pickling.
[0380] The cold rolling conditions may be the same as conventional
known conditions, and are not limited. In order to ensure
homogeneity of the structure, the reduction ratio of cold rolling
is preferably 60% or more. The reasons for limiting the conditions
of the cold-rolled sheet annealing step are described below.
[0381] The cold-rolled sheet is then subjected to the cold-rolled
sheet annealing step comprising continuous annealing and
cooling.
[0382] Continuous Annealing Temperature: Temperature in the
Ferrite-Austenite Two-Phase Coexistence Region
[0383] By annealing at a temperature in the ferrite-austenite
two-phase coexistence region, the acicular ferrite phase is formed.
In addition, the (111) aggregation structure is strongly developed
in the ferrite phase to obtain a high r value. On the other hand,
with a high temperature where an austenite single phase is formed
beyond the ferrite-austenite two-phase coexistence region, the
aggregation structure of the steel sheet is made random by reverse
transformation and transformation to decrease the r value.
Therefore, in the present invention, the annealing temperature of
continuous annealing is limited to the recrystallization
temperature or more in the ferrite-austenite two-phase coexistence
region. From the viewpoint of stability of the r value, the
temperature is preferably set so that the fraction of austenite is
10% to 50%. With the continuous annealing temperature of less than
the recrystallization temperature, ductility deteriorates, and the
steel sheet can be applied as an automobile part only to limited
special applications. Therefore, the annealing temperature is
preferably the recrystallization temperature or more.
[0384] The retention time of continuous annealing is preferably as
short as possible in order to ensure the production efficiency, the
fine structure and the amount of dissolved N. From the viewpoint of
stability of the operation, the retention time is preferably 10
seconds or more. Also, in order to ensure the fine structure and
the amount of dissolved N, the retention time is preferably 90
seconds or less. From the viewpoint of stability of material
properties, the retention time is preferably 20 seconds or
more.
[0385] Cooling After Continuous Annealing: Cooling to the
Temperature Region of 500.degree. C. or Less at a Cooling Rate of
10 to 300.degree. C./s
[0386] Cooling after soaking by continuous annealing is important
for making fine the structure, forming the acicular ferrite phase,
and ensuring the amount of dissolved N. In the present invention,
cooling is continuously carried out to the temperature region of at
least 500.degree. C. or less at a cooling rate of 10.degree. C./s
or more. With a cooling rate of less than 10.degree. C./s, a
necessary amount of acicular ferrite phase, a homogeneous fine
structure and a sufficient amount of dissolved N cannot be
obtained. On the other hand, with a cooling rate of over
300.degree. C./s, homogeneity in material properties of the steel
sheet in the width direction is insufficient. When the stop
temperature of cooling at a cooling rate of 10 to 300.degree. C./s
after continuous annealing exceeds 500.degree. C., refinement of
the structure cannot be attained.
[0387] Temper Rolling or Lever Processing: Elongation of 0.5 to
10%
[0388] In the present invention, in order to correct the shape and
control roughness, temper rolling or leveler processing may be
carried out subsequent to the cold rolling step. When the total
elongation of temper rolling or leveler processing is less than
0.5%, the desired purpose of correcting the shape and controlling
roughness cannot be achieved. On the other hand, with a total
elongation of over 10%, ductility deteriorates. In order to ensure
ductility, the elongation is preferably 5% or less. It is confirmed
that the processing system of temper rolling is different from that
of leveler processing, but the effects of both processes are
substantially the same. Temper rolling and leveler processing are
effective even after plating.
[0389] Description will be now be made of the reasons for limiting
the composition of a high-tensile-strength steel sheet according to
a seventh embodiment of the present invention.
[0390] C: 0.025 to 0.15%
[0391] C is an element for increasing the strength of a steel
sheet, and 0.025% or more of C must be contained for controlling
the structure to a homogeneous fine structure, which is an
important requirement of the present invention, and ensuring a
sufficient amount of a martensite phase. With a C content of over
0.15%, the ratio of the carbide in the steel sheet is excessively
increased to significantly deteriorate ductility and moldability.
With a C content of over 0.15%, there is the more important problem
of significantly deteriorating spot weldability and arc
weldability. Therefore, the C content is limited in the range of
0.025 to 0.15%. From the viewpoint of improvement in moldability,
the C content is preferably 0.08% or less. Particularly, when good
ductility is required, the C content is preferably 0.05% or
less.
[0392] Si: 1.0% or Less
[0393] Si is a useful component capable of increasing the strength
of the steel sheet without significantly deteriorating ductility of
steel. The Si content is preferably 0.005% or more, and more
preferably 0.10% or more. On the other hand, Si is an element which
greatly changes the transformation point during hot rolling to
cause difficulties in ensuring quality and the shape, or adversely
affects surface properties, chemical conversion properties, and the
like, particularly the beauty of the surface of the steel sheet,
and adversely affects plating properties. In the present invention,
therefore, the Si content is limited to 1.0% or less. However, the
above-described adverse effects can be kept down as long as Si is
1.0% or less. Particularly, in applications required for the steel
sheet to have a low level of strength and, particularly, surface
beauty, Si is preferably 0.5% or less.
[0394] Mn: 2.0% or Less
[0395] Mn is an element effective to prevent hot cracking with S,
and Mn is preferably added according to the amount of S contained.
Mn also has the great effect of making fine crystal grains, and is
preferably added for improving material properties. Furthermore, Mn
is an element effective to stably form martensite during rapid
cooling after continuous annealing. In order to stably fix S, the
Mn content is preferably 0.2% or more. Mn is also an element for
increasing the strength of the steel sheet, and is preferably added
in an amount of 1.2% or more when a strength TS of over 500 MPa is
required. The Mn content is more preferably 1.5% or more.
[0396] With the Mn content increased to this level, there is the
advantage that variations in the mechanical properties of the steel
sheet with respect to variations in the hot-rolling conditions,
particularly strain age hardenability, are significantly improved.
However, with the excessively high Mn content of over 2.0%, a high
r value, which is an important requirement of the present
invention, cannot be easily obtained, and ductility significantly
deteriorates. Therefore, the Mn content is limited to 2.0% or less.
In applications required to have good corrosion resistance and
moldability, the Mn content is preferably 1.7% or less.
[0397] P: 0.08% or Less
[0398] P is a useful element as a solid solution strengthening
element for steel, and is preferably added in an amount of 0.001%
or more, and more preferably 0.015% or more, from the viewpoint of
an increase in strength. On the other hand, when P is excessively
added, steel is embrittled, and stretch-flanging properties of the
steel sheet deteriorate. Also, P is liable to strongly segregate in
steel, thereby causing embrittlement of a weld. Therefore, P is
limited to 0.08% or less. In applications in which elongated flange
processability and weld toughness are considered as important, P is
preferably 0.04% or less.
[0399] S: 0.02% or Less
[0400] S is present as an inclusion in the steel sheet, decreases
ductility of the steel sheet, and causes deterioration in corrosion
resistance. Therefore, the Si content is as low as possible, and in
the present invention, the S content is limited to 0.02% or less.
Particularly, in applications required to have good processability,
S is preferably 0.015% or less. Particularly, in applications
required to have excellent stretch-flanging properties, S is
preferably 0.008% or less. Although the detailed mechanism is not
known, in order to stably maintain the strain age hardenability of
the steel sheet in a high level, it is effective to decrease the S
content to 0.008% or less.
[0401] Al: 0.02% or Less
[0402] Al is an element functioning as a deoxidization for
improving cleanliness of steel, and making fine the structure of
the steel sheet. In the present invention, the Al content is
preferably 0.001% or more. In the present invention, dissolved N is
used as a strengthening element, but aluminum killed steel
containing Al in a suitable range has mechanical properties
superior to those of conventional rimmed steel not containing Al.
On the other hand, with an excessively high Al content, the surface
properties of the steel sheet deteriorate, and the amount of
dissolved N is significantly decreased to cause difficulties in
obtaining a large amount of strain age hardening, which is the main
object of the present invention. Therefore, in the present
invention, Al is limited to 0.02% or less. From the viewpoint of
stability of material properties, Al is more preferably 0.001 to
0.015%. Although a decrease in the Al content possibly causes
coarsening of crystal grains, in the present invention, the amounts
of other alloy elements are appropriately determined to
appropriately set the annealing conditions, thereby effective
preventing
[0403] coarsening.
[0404] N: 0.0050 to 0.0250%
[0405] N is an element for increasing the strength of the steel
sheet by solid solution strengthening and strain age hardening, and
in the present invention, N is the most important element. In the
present invention, an appropriate amount of N is contained, the Al
content is controlled to the appropriate value, and production
conditions such as the hot-rolling conditions, and the annealing
conditions are controlled to ensure necessary and sufficient
dissolved N in a cold-rolled product or a coated product. This
exhibits the sufficient effect of increasing strength (yield stress
and tensile strength) by solid solution strengthening and strain
age hardening, to stably obtain the target values of the mechanical
properties of the steel sheet of the present invention, such as a
tensile strength of 440 MPa or more, an amount (BH amount) of
bake-hardening of 80 MPa or more, and an increase is tensile
strength .DELTA.TS of 40 MPa or more after strain aging. Since N
also has the function to decrease the transformation point, N is
effective for rolling of a thin material for which rolling at a
temperature greatly over the transformation point is
undesirable.
[0406] With a N content of less than 0.0050%, the effect of
increasing strength is less stably exhibited, while with a N
content of over 0.0250%, the rate of occurrence of internal defects
in the steel sheet is increased, and slab cracking frequently
occurs during continuous casting. Therefore, N is limited to the
range of 0.0050 to 0.0250%. From the viewpoint of improvement in
stability of material properties and yield over the entire
production process, N is preferably in the range of 0.0070 to
0.0170%. With the N amount in the range of the present invention,
there is no adverse effect on weldability, and the like.
[0407] Dissolved N: 0.0010% or More
[0408] In order to ensure sufficient strength of a cold-rolled
product, and effectively exhibit strain age hardening with N, it is
necessary that the content of dissolved N (solid solution N) in the
steel sheet is at least 0.0010% or more.
[0409] The amount of dissolved N is determined by subtracting the
amount of precipitated N from the total N amount of steel. As a
result of comparison research of various methods, the inventors
found that electrolytic extraction analysis using
constant-potential electrolysis is effective as the method of
analyzing the amount of precipitated N. As the method of dissolving
ferrite used for extraction analysis, an acid digestion method, a
halogen method, or an electrolysis method can be used. Of these
methods, the electrolysis method can stably dissolve only ferrite
without decomposing very unstable precipitates such as a carbide, a
nitride, etc. As the electrolyte, an acetyl-acetone system is used
for electrolysis at a constant potential. In the present invention,
the results of measurement of the amount of precipitated N by
constant-potential electrolysis showed best correspondence with
changes in actual material properties.
[0410] Therefore, in the present invention, the residue after
extraction by constant-potential electrolysis is chemically
analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated
N.
[0411] In order to obtain higher BH and .DELTA.TS, the amount of
dissolved N is preferably 0.0020% or more, more preferably 0.0020%
or more. In order to obtain further high values, the amount of
dissolved N is preferably 0.0030% or more. Although the upper limit
of the amount of dissolved N is not limited, the mechanical
properties less deteriorate even when the all amount of N added
remains.
[0412] N/Al (the Content Ratio of N to Al): 0.3 or More
[0413] In order to cause 0.0010% or more of dissolved N to stably
remain in a product state, the amount of Al which is an element for
strongly fixing N, must be limited. As a result of research of
steel sheets in which the combination of the N content (0.0050 to
0.0250%) and the Al content (0.02% or less) were widely changed in
the composition range of the present invention, it was found that
with N/Al of 0.3 or more, the amount of dissolved N of a
cold-rolled product or coated 1 product can be stably set to
0.0010% or more. Therefore, N/Al is limited to 0.3 or more.
[0414] In the present invention, the above component preferably
further contains at least one of the following groups d to g:
[0415] Group d: at least one of Cu, Ni, Cr and Mo in a total of
1.0% or less;
[0416] Group e: at least one of Nb, Ti and V in a total of 0.1% or
less;
[0417] Group f: 0.00305 of B; and
[0418] Group g: one or both of Ca and REM in a total of 0.0010 to
0.010%.
[0419] Elements of group d: Cu, Ni, Cr and Mo are all contribute to
an increase in strength of the steel sheet, and can be contained
singly or in a combination according to demand. The effect is
recognized by containing 0.005% or more each of Cu, Ni, Cr and Mo.
However, with an excessively high content, deformation resistance
in hot rolling at elevated temperatures is increased, or chemical
conversion properties and surface treatment properties in a wide
sense deteriorate, and a welded portion is hardened to deteriorate
weld moldability. Also, the r value is liable to decrease.
Therefore, the elements of group a are preferably contained in a
total of 1.0% or less. With a Mo content of 0.05% or more, the r
value is significantly decreased in some cases. In the present
invention, therefore, the Mo content is preferably limited to less
than 0.05%.
[0420] Elements of group e: All of Nb, Ti and V are elements
contributing-refinement and homegenization of crystal grains, and
may be added singly or in a combination according to demand. The
effect can be recognized by containing 0.005% or more each of Nb,
Ti and V. However, with an excessively high content, deformation
resistance in hot rolling at elevated temperatures is increased, or
chemical conversion properties and surface treatment properties in
a wide sense deteriorate. Therefore, the elements in group b are
preferably contained at a total of 0.1% or less.
[0421] Elements of group f: B is an element having the effect of
improving hardenability of steel, and can be contained for
increasing the fraction of a low-temperature transformation phase
other than the ferrite phase to increase strength of steel
according to demand. This effect is recognized with a B content of
0.0005% or more. However, with an excessively high B content,
deformability at elevated temperatures in hot rolling deteriorates
to produce BN, decreasing the amount of dissolved N. Therefore, the
B content is preferably 0.0030% or less.
[0422] Elements in group g: Both Ca and REM are elements useful for
controlling the form of inclusions. Particularly, when the stretch
flanging property is required, these elements are preferably added
singly or in a combination. When the total of the elements of group
d is less than 0.0010%, the effect of controlling the form of
inclusions is insufficient, while when the total exceeds 0.010%,
surface defects significantly occur. Therefore, the total of the
elements of group d is preferably limited to the range of 0.0010 to
0.010%. This permits improvement in the stretch flanging property
without causing surface defects.
[0423] The structure of the steel sheet of the present invention is
described below.
[0424] Area Ratio of Ferrite Phase: 80% or More
[0425] The cold-rolled steel sheet of the present invention is
directed to use as an automobile steel sheet required to have some
extent of moldability, and has a structure containing the ferrite
phase at an area ratio of 80% or more in order to ensure ductility.
With the ferrite phase at an area ratio of less than 80%, it is
difficult to ensure ductility required for an automobile steel
sheet required to have moldability. When good ductility is
required, the area ratio of the ferrite phase is preferably 85% or
more. In the present invention, "ferrite" means so-called polygonal
ferrite in which no strain remains.
[0426] Average Crystal Grain Diameter of Ferrite Phase: 10 .mu.m or
Less
[0427] In the present invention, the value used as the average
crystal grain diameter is a higher value of the value calculated
from a photograph of a sectional structure by a quadrature method
defined by ASTM, and the nominal value determined by an intercept
method defined by ASTM (refer to, for example, Umemoto et al.: Heat
Treatment, 24 (1984), p334).
[0428] The cold-rolled steel sheet of the present invention
maintains a predetermined amount of dissolved N in the product
step. However, as a result of experiment and research conducted by
the inventors, it was found that variations in strain age
hardenability occur in steel sheets containing the same amount of
dissolved N, and one of the main causes of the variations is a
crystal grain diameter. In the structure of the present invention,
in order to stably obtain a high BH amount and .DELTA.TS, the
average crystal grain diameter is at Least 10 .mu.m or less, and
preferably 8 .mu.m or less. Although the detailed mechanism is not
known, this is supposed to be related to the segregation and
precipitation of alloy elements in the crystal grain boundaries,
and the influences of processing and heat history on the
segregation and precipitation.
[0429] Therefore, in order to achieve stability of strain age
hardenability, the average crystal grain diameter of the ferrite
phase is 10 .mu.m or less, and preferably 8 .mu.m or less.
[0430] In order to ensure ductility of an automobile steel sheet,
and stability of strain age hardenability, the structure of the
present invention contains the ferrite phase with an average
crystal grain diameter of 10 .mu.m or less at an area ratio of 80%
or more.
[0431] Area Ratio of Martensite Phase: 2% or More
[0432] The cold-rolled steel sheet of the present invention
contains the martensite phase as a second phase at an area ratio of
2% or more. The presence of 2% or more of the martensite phase can
produce good ductility and a large amount of strain age hardening.
Although the detailed mechanism is not known, this effect is
supposed to be due to the effective accumulation of strain in the
steel sheet due to the presence of the martensite phase during
pre-strain processing before aging. Furthermore, the presence of
the martensite phase is effective to improve aging deterioration.
In order to a good balance between strength and ductility and a low
yield ratio, the area ratio of the martensite phase is preferably
5% or more. The presence of the martensite phase at an area ratio
of over 20% causes the problem of deteriorating ductility.
Therefore, the area ratio of the martensite phase is 2% or more,
and preferably 5% to 20%.
[0433] Besides the above-described martensite phase, pearlite,
bainite, residual austenite are present as second phases without
causing any problem. However, in the present invention, it is
necessary that the fraction of the ferrite phase is 80% or more,
and the fraction of the martensite phase is 2% or more. Therefore,
the total area ratios of pearlite, bainiate and residual austenite
are limited to less than L8%.
[0434] The cold-rolled steel sheet of the present invention, which
has the above-described composition and structure, has a tensile
strength (TS) of 440 MPa to about 780 MPa, a high r value of 1.2 or
more obtained by controlling the aggregation structure of the
ferrite base phase, and excellent strain age 5 hardenability. A
steel sheet having TS of less than 440 MPa cannot be widely applied
to members having structural components. Furthermore, in order to
extend the application range, TS is preferably 500 MPa or more.
With the r value of less than 1.2, the steel sheet cannot be
applied to a wide range of press forming parts. The preferable
range of the r value is 1.4 or more.
[0435] As described above, in the present invention, "excellent
strain age hardenability" means that in aging under conditions of
holding at a temperature of 170.degree. C. for 20 min. after
pre-deformation with a tensile strain of 5%, the increment in
deformation stress (represented by the amount of BH=yield stress
after aging-pre-deformation stress before aging) after aging is 80
MPa or more, and the increment in tensile strength (represented by
.DELTA.TS=tensile strength after aging-tensile strength without
strain aging) after strain aging (pre-deformation+aging) is 40 MPa
or more.
[0436] In defining the strain age hardenability, the amount of
pre-strain (pre-deformation) is an important factor. As a result of
research of the influence of the amount of pre-strain on strain age
hardenability, the inventors found that (1) the deformation stress
in the above-described deformation system can be referred to as an
amount of approximately uniaxial strain (tensile strain) except the
case of excessive deep drawing, (2) the amount of uniaxial strain
of an actual part exceeds 5%, and (3) the strength of a part
sufficiently corresponds to the strength (YS and TS) obtained after
strain aging with a pre-strain of 5%. In the present invention,
based on these findings, the pre-deformation of strain aging is
defined to a tensile strain of 5%.
[0437] Conventional coating and baking conditions include
170.degree. C. and 20 min as standards. When a strain of 5% is
applied to the steel sheet of the present invention, which contains
a large amount of dissolved N, hardening can be achieved even by
aging at low temperature. In other words, the range of aging
conditions can be widened. In order to attain a sufficient amount
of hardening, generally, retention at a higher temperature for a
longer time is advantageous as long as softening does not occurs by
over aging.
[0438] Specifically, in the steel sheet of the present invention,
the lower limit of the heating temperature at which hardening
significantly takes place after pre-deformation is about
100.degree. C. On the other hand, with the heating temperature of
over 300.degree. C., hardening peaks, thereby causing the tendency
to soften and significantly causing thermal strain and temper
color. With the retention time of about 30 seconds or more,
hardening can be sufficiently achieved at a heating temperature of
about 200.degree. C. In order to obtain more stable hardening, the
retention time is preferably 60 seconds or more. However, retention
for over 20 mines is practically disadvantageous because further
hardening cannot be expected, and the production efficiency
significantly deteriorates.
[0439] Therefore, in the present invention, the conventional
coating and baking conditions, i.e., the heating temperature of
170.degree. C. and the retention time of 20 minutes, are set as the
aging conditions. With the steel sheet of the present invention,
hardening can be stably achieved even under the aging conditions of
a low heating temperature and a short retention time, which fail to
achieve sufficient hardening in a conventional bake-hardening steel
sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and
other methods such as induction heating, heating with a
nonoxidation flame, a laser, plasma, or the like, etc. can be
preferably used.
[0440] The strength of an automobile part must be sufficient to
resist an external complicated stress load, and thus not only
strength in a low strain region but also strength in a high strain
region are important for a raw material steel sheet. In
consideration of this point, in the steel sheet of the present
invention used as a raw material for automobile parts, BH is 80 MPa
or more, and .DELTA.TS is 40 MPa or more. More preferably, BH is
100 MPa or more, and .DELTA.TS is 50 MPa or more. In order to
further increase BH and TS, the heating temperature in aging may be
set to a higher temperature, and/or the retention time may be set
to a longer time.
[0441] The steel sheet of the present invention has the advantage
that when the steel sheet is allowed to stand at room temperature
for about one week without heating after forming, an increase in
strength of about 40% of that at the time of complete aging can be
expected.
[0442] The steel sheet of the present invention also has the
advantage that even when it is allowed in an unmolded state at room
temperature for a long time, aging deterioration (an increase in YS
and a decrease in El (elongation)) does not occurs, unlike a
conventional aging steel sheet. In order to prevent the occurrence
of a trouble in actual press forming, it is necessary that in aging
at room temperature for 3 months before press forming, an increase
in YS is 30 MPa or less, a decrease in elongation is 2% or less,
and a recovery of yield point elongation is 0.2% or less.
[0443] In the present invention, the surface of the cold-rolled
steel sheet may be coated by hot-dip galvanization or alloying
hot-dip galvanization without any problem, and TS, BH and .DELTA.TS
are equivalent to those before plating. AS the plating method,
electro-galvanization, hot-dip galvanization, alloying hot-dip
galvanization, electro-tinning, electric chromium plating,
electro-nickeling, and the like may be preferably used.
[0444] The method of producing a steel sheet according to an eighth
embodiment of the present invention will be described.
[0445] The steel sheet of the present invention is basically
produced by performing the hot rolling step in which a steel slab
having the above-described composition is heated, and then roughly
rolled to form a sheet bar, and the sheet bar is finish-rolled and
cooled to form a coiled hot-rolled sheet, the cold rolling step in
which the hot-rolled sheet is pickled and cold-rolled to form a
cold-rolled sheet, and the cold-rolled sheet annealing step in
which the cold-rolled sheet is box-annealing and then continuously
annealed.
[0446] Although the slab used in the production method of the
present invention is preferably produced by a continuous casting
method in order to prevent macro-segregation of components, an
ingot making method or a thin slab casing method may be used.
Alternatively, a conventional method comprising cooling the
produced slab to room temperature and then again heating the slab,
or an energy-saving process of direct rolling comprising charging a
hot slab into a heating furnace without cooling and then rolling
it, or rolling directly the slab immediately after slightly keeping
it warm may be used without a problem. Particularly, direct rolling
is a useful technique for effectively ensuring dissolved N.
[0447] Description will be now be made of the reasons for limiting
the conditions of the hot rolling step.
[0448] Slab Heating Temperature: 1000.degree. C. or More
[0449] The slab heating temperature is preferably 1000.degree. C.
or more in order to ensure a necessary and sufficient amount of
dissolved N in an initial state and satisfy the target amount of
dissolved N in a product. Since a loss is increased by an increase
in the oxide weight, the heating temperature is preferably
1280.degree. C. or less.
[0450] The slab heated under the above condition is roughly rolled
to form a sheet bar. The condition of the rough rolling is not
defined, and rough rolling may be performed according to a normal
method. However, in order to ensure the amount of dissolved N,
rough rolling is preferably performed within a as short time as
possible. Then, the sheet bar is finish-rolled to form a hot-rolled
sheet.
[0451] In the present invention, the adjacent sheet bars are
preferably bonded together during the time between rough rolling
and finish rolling, and then continuously rolled. As the bonding
means, a pressure welding method, a laser welding method, an
electron beam welding method, or the like is preferably used.
[0452] By continuous rolling, so-called non-stationary portions at
the front end and the rear end of a coil (processed material) are
removed to permit hot rolling over the entire length and the entire
width of the coil (processed material) under stable conditions.
This is very effective to improve the sectional shape and
dimensions of not only the hot-rolled steel sheet but also the
cold-rolled steel sheet. When the steel sheet is cooled on a hot
run table after rolling, the shape of the steel sheet can be
sufficiently maintained because tension can be always applied.
[0453] By continuous rolling, the ends of the coil can be stably
passed, and it is thus possible to use lubricating rolling, which
cannot be easily applied to ordinary single rolling for each sheet
bar because of the problem of continuous rolling processes and
biting property. Therefore, the rolling load can be decreased, and
at the same time, the surface pressure of the roll can be
decreased, thereby increasing the life time of the roll.
[0454] In the present invention, at the entrance of a finisher
between rough rolling and finish rolling, the temperature
distributions of the sheet bar in the width direction and the long
direction thereof are preferably made uniform by using any one or
both of a sheet bar edge heater for heating the ends of the sheet
bar in the width direction and a sheet bar heater for heating the
ends of the sheet bar in the long direction. This can further
decrease the variations in material properties of the steel sheet.
The sheet bar edge heater and the sheet bar heater are preferably
of an induction heating type.
[0455] As the procedure of use of the heaters, preferably, the
difference in temperature in the width direction is first corrected
by the sheet bar edge heater. Although the heating amount depends
upon the steel composition, the temperature distribution in the
width direction at the finisher entrance is preferably set in the
range of about 20.degree. C. or less. Next, the difference in
temperature in the long direction is corrected by the sheet bar
heater. At this time, the heating amount is preferably set so that
the temperatures at the ends in the long direction are about
20.degree. C. higher than the temperature at the center.
[0456] Finisher Delivery Temperature: 800.degree. C. or More
[0457] In order to obtain a homogeneous fine hot-rolled base sheet
structure, the finisher deliver temperature FDT is 800.degree. C.
or more. With a FDT of lower than 800.degree. C., the structure of
the steel sheet becomes inhomogeneous, and the processed structure
partially remains to leave heterogeneity of the structure after the
cold-rolled sheet annealing step. Therefore, the danger of causing
various troubles in press forming is increased. When a high coiling
temperature is used for avoiding the processed structure from
remaining, coarse crystal grains are produced to cause the same
troubles as described above. With the high coiling temperature, the
amount of dissolved N is significantly decreased to cause
difficulties in obtaining a target tensile strength of 440 MPa or
more. Therefore, the finisher deliver temperature FDT is
800.degree. C. or more. In order to further improve the mechanical
properties, the FDT is preferably 820.degree. C. or more. Although
the upper limit of FDT is not limited, a scale scar significantly
occurs at excessively high FDT. The FDT is preferably up to about
1000.degree. C.
[0458] Although cooling after finish rolling is not strictly
limited, the conditions described below are preferable from the
viewpoint of homogeneity in material properties of the steel sheet
in the long direction and the width direction thereof. In the
present invention, cooling is preferably started immediately after
(within 0.5 seconds after) finish rolling, and the mean cooling
rate in cooling is preferably 40.degree. C./s or more. By
satisfying these conditions, the steel sheet can be rapidly cooled
in the high temperature region where AlN precipitates to
effectively ensure N in a solid solution state. When the starting
time of cooling or the cooling rate does not satisfy the above
condition, grain growth excessively proceeds to fail to achieve
fine crystal grains, and promote AlN precipitation due to stain
energy introduced by rolling. Therefore, the amount of dissolved N
tends to decrease, and the structure tends to be made
inhomogeneous. In order to ensure homogeneity in material
properties and shape, the cooling rate is preferably kept at
300.degree. C./s or less.
[0459] Coiling Temperature: 800.degree. C. or Less
[0460] The strength of the steel sheet is liable to increase as the
coiling temperature CT decreases. In order to ensure the target
tensile strength TS of 440 MPa or more, the CT is preferably
800.degree. C. or less. With a CT of less than 200.degree. C., the
shape of the steel sheet is readily disturbed to increase the
danger of causing troubles in a practical operation, thereby
deteriorating homogeneity of material properties. Therefore, the CT
is preferably 200.degree. C. or more. When the homogeneity of the
material properties is required, the CT is preferably 300.degree.
C. or more, and more preferably 350.degree. C. or more. In the
present invention, in finish rolling, lubricating rolling may be
performed for decreasing the hot rolling load. The lubricating
rolling has the effect of further making homogeneous the shape and
material properties of the hot-rolled sheet. During the lubricating
rolling, the frictional coefficient is preferably in the range of
0.25 to 0.10. By combining lubricating rolling and continuous
rolling, the operation of hot rolling is further stabilized.
[0461] The hot-rolled sheet subjected to the above-described hot
rolling step is then pickled and cold-rolled in the cold rolling
step to form a cold-rolled sheet.
[0462] The pickling conditions may be the same as conventional
known conditions, and are not limited. When the scale of the
hot-rolled sheet is extremely small, cold rolling may be
immediately after hot rolling without pickling.
[0463] The cold rolling conditions may be the same as conventional
known conditions, and are not limited. In order to ensure
homogeneity of the structure, the reduction ratio of cold rolling
is preferably 40% or more. The reasons for limiting the conditions
of the cold rolling step are described below.
[0464] The cold-rolled sheet is then subjected to the cold-rolled
sheet annealing step comprising box annealing and continuous
annealing.
[0465] Box Annealing Temperature: the Recrystallization Temperature
to 800.degree. C.
[0466] In the present invention, the cold-rolled sheet is subjected
to box annealing to control the aggregation structure of the
ferrite phase as a base. By controlling the aggregation structure
of the ferrite phase, the r value of the produced sheet can be
increased. By box annealing, the (111) aggregation structure
suitable for increasing the r value is readily formed in the
produced sheet.
[0467] With the box annealing temperature less than the
recrystallization temperature, recrystallization is not completed
to fail to control the aggregation structure of the ferrite phase,
thereby failing to increase the r value. On the other hand, with
the box annealing temperature of over 800.degree. C., surface
defects significantly occur in the steel sheet, thereby failing to
achieve the initial purpose. Box annealing is preferably performed
in an annealing atmosphere containing a nitrogen gas as a main
component and 3 to 5% of hydrogen gas. In this case, the heating
and cooling rates may be the same as normal box annealing, and are
about 30.degree. C./hr. By using 100% hydrogen gas as an annealing
atmosphere gas, the higher heating and cooling rates may be
used.
[0468] Continuous Annealing Temperature: Ac.sub.1 Transformation
Point to (Ac.sub.3 Transformation Point--20.degree. C.)
[0469] With the continuous annealing temperature of less than the
Ac.sub.1 transformation point, the martensite phase is not formed
after annealing, while with the continuous annealing temperature of
over (Ac.sub.3 transformation point--20.degree. C.), the desirable
aggregation structure formed in box annealing is lost due to
transformation, thereby failing to obtain the produced sheet having
a high r value. Therefore, the continuous annealing temperature is
preferably Ac.sub.1 transformation point to (Ac.sub.3
transformation point--20.degree. C.). The retention time of
continuous annealing is preferably as short as possible in order to
ensure the production efficiency, the fine structure and the amount
of dissolved N. From the viewpoint of stability of the operation,
the retention time is preferably 10 seconds or more. Also, in order
to ensure the fine structure and the amount of dissolved N, the
retention time is preferably 120 seconds or less. From the
viewpoint of stability of material properties, the retention time
is preferably 20 seconds or more.
[0470] Cooling After Continuous Annealing: Cooling to the
Temperature Region of 500.degree. C. or Less at a Cooling Rate of
10 to 300.degree. C./s
[0471] Cooling after soaking by continuous annealing is important
for making fine the structure, forming the martensite phase, and
ensuring the amount of dissolved N. In the present invention,
cooling is continuously carried out to the temperature region of at
least 500.degree. C. or less at a cooling rate of 10.degree. C./s
or more. With a cooling rate of less than 10.degree. C./s, a
necessary amount of martensite phase, a homogeneous fine structure
and a sufficient amount of dissolved N cannot be obtained. On the
other hand, with a cooling rate of over 300.degree. C./s,
homogeneity in material properties of the steel sheet in the width
direction deteriorates due to a significant increase in the amount
of supersaturated dissolved C. When the stop temperature of cooling
at a cooling rate of 10 to 300.degree. C./s after continuous
annealing exceeds 500.degree. C., refinement of the structure
cannot be attained.
[0472] Over Aging Condition: Retention in the Temperature Region of
350.degree. C. to the Cooling Stop Temperature for 20 Seconds or
More Subsequent to Cooling After Continuous Annealing
[0473] Over aging may be performed by retention in the temperature
region of 350.degree. C. to the cooling stop temperature for 20
seconds or more subsequent to the stop of cooling after soaking by
continuous annealing. By over aging, the amount of dissolved C can
be selectively decreased, while the amount of dissolved N is
maintained. With the retention temperature region of less than
350.degree. C., a long time is required for decreasing the amount
of dissolved C to cause a reduction in productivity. Therefore, the
temperature region is preferably 350.degree. C. or more.
[0474] By retention in the temperature region of 350.degree. C. to
the cooling stop temperature for 20 seconds or more, the amount of
dissolved C can be decreased to achieve a higher degree of
non-aging properties at room temperature. By increasing the
retention time, further improvement can be expected, but the effect
is saturated with the retention time of about 120 seconds.
Therefore, the retention time is preferably 120 second or less.
[0475] In order to obtain a larger amount of strain age hardening,
it is advantageous to use either of dissolved C and dissolved N.
However, by using dissolved C, aging deterioration at room
temperature becomes significant, thereby limiting parts to which
the steel sheet is applied. Therefore, in order to produce a strain
age hardenable steel sheet having versatility, over aging is
preferably performed with the sufficient amount of dissolved N
being ensured.
[0476] In producing a high-tensile-strength cold-rolled coated
steel sheet comprising a high-tensile-strength cold-rolled steel
sheet and a hot-dip coated layer formed on the surface thereof,
continuous annealing after box annealing can be performed in a
continuous hot-dip coating line comprising hot-dip galvanization
subsequent to cooling after continuous annealing or further
alloying to produce a hot-dip galvanized steel sheet.
[0477] Temper Rolling or Lever Processing: Elongation of 0.2 to
15%
[0478] In the present invention, in order to correct the shape and
control roughness, temper rolling or leveler processing may be
carried out subsequent to the cold rolling step. When the total
elongation of temper rolling or leveler processing is less than
0.2%, the desired purpose of correcting the shape and controlling
roughness cannot be achieved. On the other hand, with a total
elongation of over 15%, ductility significantly deteriorates. It is
confirmed that the processing system of temper rolling is different
from that of leveler processing, but the effects of both processes
are substantially the same. Temper rolling and leveler processing
are effective after plating.
[0479] For reference, description will now be made of forming
conditions and conditions for subsequent heat treatment for
increasing strength when the steel sheet of the present invention
is molded, for example, press-molded. When the steel sheet of the
present invention is subjected to press working, for example, deep
drawing, the strain introduced by press working is several % to
several tens %. Although the amount of strain changes with molded
parts, a strain of about 5 to 10% is introduced into an inner plate
and a structural member in the automobile field.
[0480] These automobile parts are heat-treated by coating and
baking. However, with the steel sheet of the present invention,
strength of a molded product can be effectively increased after
heat treatment. In the present invention, as a method of evaluating
burning hardenability in a laboratory, a tensile test specimen of
JIS No. 5 size is obtained from the steel sheet in the rolling
direction, and tensile strain of 10% is applied to the tensile test
specimen by a tensile testing machine. Then, the specimen is
heat-treated and again subjected to a tensile test. Particularly,
when properties are evaluated after heat treatment in a low
temperature region, the heat treatment conditions include
120.degree. C. and 20 minutes. In this test, the properties of the
completed portion after heat treatment subsequent to press forming
are evaluated.
[0481] Namely, in the present invention, the difference (.DELTA.TS)
between the tensile strength of the specimen after application of
tensile strain and heat treatment and the tensile strength of a
product is defined as the strength increasing ability of heat
treatment.
[0482] In order to increase the strength of the molded product, the
amount of strain introduced by forming, or the heat treatment
temperature after processing is preferably as high as possible.
[0483] However, with the steel sheet of the present invention, when
the amount of applied strain is about 5 to 10%, the strength can be
sufficiently increased even by heat treatment at a temperature
lower than conventional heat treatment, i.e., a temperature of
200.degree. C. or less, after forming. However, with a heat
treatment temperature of less than 120.degree. C., the strength
cannot be sufficiently increased with the low train applied. On the
other hand, with the heat treatment temperature of over 350.degree.
C. after forming, softening proceeds. Therefore, the temperature of
heat treatment after forming is preferably about 120 to 350.degree.
C.
[0484] The heating method is not limited, and hot gas heating,
infrared furnace heating, hot-bath heating, direct current heating,
induction heating, and the like can be used. Alternatively, only a
portion where strength is desired to be increased is selectively
heated.
EXAMPLES
[0485] In the examples below, the amount of dissolved N, the
microstructure, tensile properties, the r value, strain age 15
hardenability, and aging property were examined. The examination
methods were as follows:
[0486] (1) Amount of dissolved N
[0487] The amount of dissolved N was determined by subtracting the
amount of precipitated N from the total N amount of steel
determined by chemical analysis. The amount of precipitated N was
determined by an analysis method using a constant-potential
electrolytic method.
[0488] (2) Microstructure
[0489] A test specimen was obtained from each of cold-rolled
annealed steel sheets, and the microstructure of a section (C
section) perpendicular to the rolling direction was imaged with an
optical microscope or a scanning electron microscope. Then, the
fraction of the ferrite texture and the type and the structure
fraction of a second phase were determined by an image analysis
apparatus.
[0490] (3) Crystal Grain Diameter
[0491] In the present invention, the value used as the average
crystal grain diameter was a higher one of the value calculated
from a photograph of a sectional structure by a quadrature method
defined by ASTM, and the nominal value determined from a photograph
of a sectional structure by an intercept method defined by ASTM
(refer to, for example, Umemoto et al.: Heat Treatment, 24 (1984),
p334).
[0492] (4) Tensile Properties
[0493] A test specimen of JIS No. 5 was obtained from each of
cold-rolled annealed steel sheets in the rolling direction, and a
tensile test was carried out with a strain rate of
3.times.10.sup.-3/s according to the regulations of JIS Z 2241 to
determine yield stress YS, tensile strength TS, and elongation
El.
[0494] (5) Strain Age Hardenability
[0495] A test specimen of JIS No. 5 was obtained from each of
cold-rolled annealed steel sheets in the rolling direction, and a
tensile strain of 5% was applied as pre-deformation. Then, the
specimen was subjected to heat treatment corresponding to coating
and baking at 170.degree. C. for 20 minutes, and a tensile test
with a strain rate of 3.times.10.sup.-3/s was performed to
determine the tensile properties (yield stress TSBH, tensile
strength TSBH). Then, BH amount=YSBH-YS5%, and .DELTA.TS=TSBH-TS
were calculated. YS5% represents deformation stress in 5%
pre-deformation of the produced sheet, YSBH and TSBH represent
yield stress and tensile strength, respectively, after
pre-deformation and heat treatment, and TS represents the tensile
strength of the produced sheet.
[0496] (6) Measurement of r Value
[0497] A test specimen of JIS No. 5 was obtained from each of the
cold-rolled annealed steel sheets in each of the rolling direction
(L direction), the direction (D direction) at 45.degree. with the
rolling direction, and the direction (C direction) at 90.degree.
with the rolling direction. The width-direction strain and
thickness-direction strain of each of the test specimens were
determined when a uniaxial tensile strain of 15% was applied to
each specimen, and the r value of each specimen in each of the
directions was determined from the following ratio of
width-direction strain to thickness-direction strain:
r=ln (w/w0)/ln (t/t0)
[0498] (wherein w0 and t0 represent the width and thickness of a
specimen before the test, and w and t represent the width and
thickness of a specimen after the test).
[0499] The mean value was determined by the following equation:
rmean=(rL+2rD +rD)/4
[0500] wherein rL represents the r value in the rolling direction
(L direction), rD represents the r value in the direction (D
direction) at 45.degree. with the rolling direction, and rL
represents the r value in the direction (C direction) at 90.degree.
with the rolling direction. In order to improve the precision of
experiment, calculation was made by using changes in elongation
strain and strain in the width direction on the assumption that the
volume was constant.
[0501] (7) Aging Properties
[0502] A test specimen of JIS No. 5 was obtained from each of
cold-rolled annealed steel sheets in the rolling direction, and
then subjected to aging at 50.degree. C. for 200 hours, followed by
a tensile test. The difference in yield elongation .DELTA.Y-El
between before and after aging was determined from the obtained
results to evaluate aging properties at normal temperature. When
.DELTA.Y-El was zero, it was evaluated that the specimen has
non-aging properties and excellent natural aging resistance.
[0503] (8) Tensile Strength After Forming and Heat Treatment
[0504] A test specimen of JIS No. 5 was obtained from each of
produced sheets in the rolling direction, and then a pre-strain of
10% was applied thereto. Then, heat treatment was conducted for 20
minutes at a conventional heat treatment temperature of 120.degree.
C. and a temperature of 170.degree. C. corresponding to coating and
baking, and then tensile strength was determined.
[0505] (9) Decrease (.DELTA.El) in Total Elongation by Natural
Aging
[0506] The decrease (.DELTA.El) in total elongation by natural
aging was determined as the difference between the total elongation
measured with a specimen of JIS NO 5 obtained from the produced
sheet in the rolling direction, and the total elongation measured
with a specimen of JIS NO 5 separately obtained from the produced
sheet in the rolling direction after accelerated aging (retention
at 100.degree. C. for 8 hours) of natural aging.
Example 1
[0507] A steel slab having each of the compositions shown in Table
1 was hot-rolled to a hot-rolled sheet having a thickness of 3.5
mm, and then cold-rolled to a cold-rolled sheet having a thickness
of 0.7 mm under the conditions shown in Table 2. Then, the
cold-rolled sheet was recrystallized, annealed and further
galvannealed in a continuous annealing line or a continuous
annealing and galvanizing line. Then, the annealed sheet was
temper-rolled with a rolling reduction ratio of 1.0% to produce a
cold-rolled steel sheet and a galvannealed steel sheet having both
sides coated with a weight of 45 g/m.sup.2 per side. In Table 2,
the finisher deliver temperatures in the hot rolling process of
Steel Nos. 3 and 8 are less than the Ar.sub.3 transformation point,
and the finisher deliver temperatures of the others are the
Ar.sub.3 transformation point or more.
[0508] The thus-obtained cold-rolled steel sheets and the
galvannealed steel sheets were measured with respect to tensile
strength, the r value, and a change in tensile strength after
forming and heat treatment. The results are shown in Table 3.
[0509] Table 3 indicates that with all the cold-rolled steel sheets
and the galvannealed steel sheets obtained according to the present
invention, a high r value and excellent strain age hardenability
are obtained, as compared with comparative examples. Particularly,
in the suitable examples in which the crystal grain diameter is 20
im or less, the decrease in elongation due to natural aging is also
as low as 2.0% or less.
Example 2
[0510] A slab of steel symbol B shown in Table 1 was hot-rolled
under the same production conditions as No. 2 shown in Table 2 in
which the heating temperature was 1100.degree. C., and the finisher
deliver temperature of hot rolling was 900.degree. C., and then
coiled at coiling temperature of 550.degree. C. into a coil. The
thus-obtained coil was cold-rolled with a reduction ratio of 80%,
and then recrystallized and annealed at 840.degree. C. With respect
to the product properties of the resultant cold-rolled steel sheet,
tensile strength TS was 365 MPa, and the r value was 1.7. A test
specimen of JIS No. 5 was obtained from the cold-rolled steel sheet
in the rolling direction, and a tensile strain of 10% was applied
by a tensile test machine. Then, the specimen was subjected to heat
treatment under the heat treatment conditions (temperature and
time) shown in Table 4, and a tensile test was again performed.
Table 4 also shows the increase in tensile strength (.DELTA.TS)
from the tensile strength (TS=365 MPa) of a product before
application of strain.
[0511] Table 4 indicates that the increase in strength increases as
the heat treatment temperature increases, and the heat treatment
time increases. However, with the steel sheet of the present
invention, a sufficient increase in tensile strength of 82 MPa (85%
or more of an increase in heat treatment for 20 minutes) can be
obtained even by heat treatment at low temperature of 120.degree.
C. for a short retention time of 2 minutes. It is thus found that
with the steel sheet of the present invention, good strain age
hardenability can be obtained even by heat treatment at a low
temperature for a short time. In order to obtain the stable effect
of increasing strength of an automobile structural member, or the
like, heat treatment at a normal temperature for a normal time
causes no problem. It was confirmed that with the galvanized steel
sheets and the galvannealed steel sheets obtained by hot-dip
galvanizing and heat alloying the cold-rolled sheets, the same
results as shown in Table 4 are obtained.
Example 3
[0512] A steel slab having each of the compositions shown in Table
6 was hot-rolled under the conditions shown Table 7 to form a
hot-rolled sheet having a thickness of 3.5 mm. Each of the
thus-obtained hot-rolled sheets was cold-rolled under the
conditions shown in Table 7 to form a cold-rolled sheet having a
thickness of 0.7 mm, and then recrystallized and annealed under the
conditions shown in the same table. Some of the annealed sheets
were further coated by hot-dip galvanization or alloying hot-dip
galvanization under the conditions shown in the same table. The
thus-obtained produced sheets were examined with respect to the
amount of dissolved N, the microstructure, tensile properties, and
strain age hardenability.
[0513] The results are shown in Table 8. Table 8 indicates that
with the steel sheets of the present invention, TS.times.r
value.gtoreq.750 MPa (in a combination with at least one of B, Nb,
Ti and V, Ts.times.r value 850>MPa), BH.gtoreq.80 MPa and
.DELTA.TS.gtoreq.40 MPa are satisfied, while in the comparative
examples, at least one of the three properties does not reach the
level of the present invention. Example 4 Examples of the present
invention are described below.
[0514] Melted steel having each of the compositions shown in Table
9 was formed in an ingot by a converter, and then formed in a steel
slab by a continuous casting method. Each of the resultant steel
slabs was heated under the conditions shown in Table 10, and
roughly rolled to form a sheet bar. The sheet bar was hot-rolled by
the hot rolling step comprising finish rolling under the conditions
shown in Table 10 to form a hot-rolled rolled sheet. The Ar.sub.3
transformation point was measured by simulating the conditions for
hot rolling finish rolling using a processing transformation
measuring apparatus (produced by Fuji Denpa Kouki). The results are
shown in Table 10.
[0515] Each of the thus-obtained hot-rolled sheets was cold-rolled
by the cold rolling step under the conditions shown in Table 10 to
form a cold-rolled sheet. Then, each of the cold-rolled sheets was
continuously annealed under the conditions shown in Table 10. Some
of the cold-rolled sheets were further temper-rolled after the
cold-rolled sheet annealing step.
[0516] The thus-obtained cold-rolled annealed sheets were examined
with respect to the amount of dissolved N, the microstructure,
tensile properties, the r value, strain age hardenability and aging
properties.
[0517] The surface of each of the steel sheets of Nos. 4 and 10 was
hot-dip galvanized to form a coated steel sheet, and evaluated with
respect to the same properties as described above.
[0518] The results are shown in Table 11.
[0519] All examples of the present invention have excellent
ductility, an extremely high BH amount and .DELTA.TS, excellent
strain age hardenability, a mean r value of as high as 1.2 or more,
and non-aging properties at natural aging (excellent natural aging
resistance). With respect to the properties of the coated steel
sheets obtained by hot-dip galvanizing the surfaces of the steel
sheets of Nos. 4 and 10, the mean r value is decreased by 0.2, and
the elongation El is decreased by about 1%, as compared with the
cold-rolled steel sheets, because of shrinkage restriction of the
coated layer in the width direction. However, strain age
hardenability and natural aging resistance are substantially the
same as those before coating. On the other hand, in the comparative
examples out of the range of the present invention, ductility
deteriorates, the BH amount and .DELTA.TS are low, or natural aging
deterioration significantly occurs. Therefore, the comparative
examples do not have all the intended properties, and thus cannot
be said as steel sheets having sufficient properties.
[0520] Steel sheet No. 11 contains C, Al, N, and N/Al out of the
range of the present invention, and thus the r value, the BH
amount, .DELTA.TS and natural aging resistance deteriorate. Steel
sheet No. 12 contains B and Nb out of the range of the present
invention, and thus the amount of acicular ferrite is greatly
deviated from the range of the present invention, deteriorating the
BH amount, .DELTA.TS, and natural aging resistance.
[0521] Steel sheet No. 13 contains B out of the range of the
present invention, and thus the amount of acicular ferrite is
greatly deviated from the range of the present invention,
deteriorating the r value, the BH amount, .DELTA.TS, and natural
aging resistance. Steel sheet No. 14 contains Nb out of the range
of the present invention, and thus the amount of dissolved N is
greatly lower than the range of the present invention,
deteriorating strain age hardenability.
[0522] Steel sheet No. 15 contains N out of the range of the
present invention, and thus the amount of dissolved N is low,
deteriorating strain age hardenability. In Steel Nos. 17 to 20, the
hot rolling conditions and the cold-rolled sheet annealing
conditions are deviated from the suitable ranges, and thus the
microstructure is out of the range of the present invention,
decreasing the BH amount and .DELTA.TS and deteriorating strain age
hardenability and natural aging resistance.
Example 5
[0523] Steel having the composition shown in Table 12 was formed in
a slab by the same method as Example 4, and then heated and
temper-rolled under the conditions shown in Table 13 to form a
sheet bar. The sheet bar was then hot-rolled by the hot rolling
step comprising finish rolling under the conditions shown in Table
13 to form a hot-rolled sheet. The adjacent sheet bars on the
finisher entrance side after rough rolling were bonded together by
a melt welding method and then continuously rolled. The
temperatures of the ends of the sheet bar were controlled in the
width direction and the length direction by using an induction
heating-type sheet bar edge heater and a sheet bar heater.
[0524] The thus-obtained hot-rolled sheet was cold-rolled by the
cold rolling step comprising pickling and cold rolling under the
conditions shown in Table 13 to form a cold-rolled sheet having a
thickness of 1.6 mm. Then, the cold-rolled sheet was continuously
annealed under the conditions shown in Table 13.
[0525] The thus-obtained cold-rolled annealed sheet was examined
with respect to the amount of dissolved N, the microstructure,
tensile properties, the r value, and strain age hardenability by
the same methods as Example 4. The tensile property of each
cold-rolled annealed sheet was measured at ten positions in each of
the width direction and the long direction to examine variations in
yield strength, tensile strength and elongation.
[0526] The results are shown in Table 14.
[0527] All the examples of the present invention have excellent
strain age hardenability and a high r value, and exhibit extremely
high stable BH amount, .DELTA.TS and mean value regardless of
variations in production conditions. It was also recognized that in
the examples of the present invention, by performing continuous
rolling and controlling the temperature of the sheet bar in the
long direction and the width direction, the thickness precision and
the shape of the produced steel sheet are improved, and variations
in material properties are decreased to 1/2. Even when the
elongation of temper rolling is changed to 0.5 to 2%, and the
elongation of the leveler is changed to 0 to 1%, strain age
hardenability does not deteriorate.
Example 6
[0528] Examples of the present invention are described below.
[0529] Melted steel having each of the compositions shown in Table
15 was formed in an ingot by a converter, and then formed in a slab
by a continuous casting method. Each of the steel slabs was heated
(in some cases, a hot slab was charged) and roughly rolled under
the conditions shown in Table 16 to form a sheet bar. The sheet bar
was then hot-rolled by the hot rolling step comprising finish
rolling under the conditions shown in Table 16 to form a hot-rolled
sheet. With some of the sheet bars, the adjacent sheet bars were
bonded by the melt welding method, and then continuously
rolled.
[0530] Each of the resultant hot-rolled sheets was cold-rolled in
the cold rolling step comprising pickling and cold rolling under
the conditions shown in Table 16 to form a cold-rolled sheet. Each
of the thus-obtained cold-rolled sheet was box-annealed and then
continuously annealed under the conditions shown in Table 16. Some
of the cold-rolled sheets were temper-rolled after the cold-rolled
sheet annealing step. Box annealing may not be carried out. In all
cases, the annealing temperature of box annealing was the
recrystallization temperature or more.
[0531] The thus-obtained cold-rolled annealed sheets were examined
with respect to the amount of dissolved N, the microstructure,
tensile properties, the r value, strain age hardenability, and the
aging property.
[0532] The surfaces of the steel sheets of Nos. 17 and 18 were
coated by hot-dip galvanization in an in-line after continuous
annealing shown in Table 16 to form coated steel sheets. The coated
steel sheets were also examined with respect to the same properties
as described above.
[0533] The results are shown in Table 17.
[0534] All the examples of the present invention have excellent
ductility, extremely high stable BH amount and .DELTA.TS, excellent
strain age hardenability, a mean r value of as high as 1.2, and
natural non-aging properties. The properties of the hot-dip
galvanized steel sheets of Nos. 17 and 18 shown in Table 17 are
substantially the same as the cold-rolled steel sheets subjected to
continuous annealing. On the other hand, in the comparative
examples out of the range of the present invention, ductility
deteriorates, the BH amount and TS are low, or aging deterioration
significantly occurs. Therefore, the comparative examples do not
have all the intended properties, and thus cannot be said as steel
sheets having sufficient properties.
[0535] Steel sheet No. 11 contains C and N in amounts out of the
range of the present invention, and has an amount of dissolved N
and a martensite amount lower than the range of the present
invention. Therefore, the BH amount and .DELTA.TS are decreased,
and AY-El is increased. Steel sheet No. 12 contains Al, N/Al and N
out of the range of the present invention, and has an amount of
dissolved N lower than the range of the present invention, and the
average crystal grain diameter of ferrite higher than the range of
the present invention. Therefore, the BH amount and .DELTA.TS are
decreased, and .DELTA.Y-El is increased.
[0536] In steel sheet No. 13, the slab heating temperature and
finisher delivery temperature FDT are out of the range of the
present invention, the amount of dissolved N and the amount of
martensite are lower than the range of the present invention, and
the average crystal grain diameter of ferrite is higher than the
range of the present invention. Therefore, the r value, the BH
amount and .DELTA.TS are decreased. In steel sheet No. 14, the
coiling temperature after hot rolling is out of the range of the
present invention, the amount of dissolved N and the amount of
martensite are lower than the range of the present invention, and
the average crystal grain diameter of ferrite is higher than the
range of the present invention, Therefore, the r value, the BH
amount and .DELTA.TS are decreased.
[0537] In steel sheet No. 15, the continuous annealing temperature
is out of the range of the present invention, martensite is not
formed, and the average crystal grain diameter of ferrite is higher
than the range of the present invention. Therefore, the BH amount
and .DELTA.TS are decreased, and .DELTA.Y-El is increased. In steel
sheet No. 16, box annealing is not performed to fail to develop the
desirable aggregation structure, deteriorating the r value. Also,
the average crystal grain diameter of ferrite, and the area ratio
of martensite are out of the range of the present invention.
Example 7
[0538] Steel having the composition shown in Table 18 was formed in
a slab by the same method as Example 1, and then heated and roughly
rolled under the conditions shown in Table 19 to form a sheet bar
having a thickness of 30 mm. The sheet bar was hot-rolled by the
hot rolling step comprising finish rolling under the conditions
shown in Table 19 to form a hot-rolled sheet. With some of the
sheet bars, the adjacent sheet bars on the finisher entrance side
after rough rolling were bonded together by the melt welding
method, and then continuously rolled. The temperatures of the ends
of the sheet bar were controlled in the width direction and the
length direction by using an induction heating-type sheet bar edge
heater and a sheet bar heater.
[0539] The thus-obtained hot-rolled sheet was cold-rolled by the
cold rolling step comprising pickling and cold rolling under the
conditions shown in Table 19 to form a cold-rolled sheet having a
thickness of 1.6 mm. Then, the cold-rolled sheet was box-annealed
and then continuously annealed by a continuous annealing furnace
under the conditions shown in Table 19. In all cases, the annealing
temperature of box annealing are the recrystallization temperature
or more.
[0540] The thus-obtained cold-rolled annealed sheet was examined
with respect to the amount of dissolved N, the microstructure,
tensile properties, the r value, and strain age hardenability by
the same methods as Example 1. The tensile property of each
cold-rolled annealed sheet was measured at ten positions in each of
the width direction and the long direction to examine variations in
yield strength, tensile strength and elongation. The variation is
shown by a difference between the maxim and minimum of all
measurements, for example, .delta.YS=(maximum of YS)-(minimum of
YS). The results are shown in Table 20.
[0541] All the examples of the present invention have excellent
strain age hardenability and a high r value, and exhibit extremely
high stable BH amount, .DELTA.TS and mean r value regardless of
variations in production conditions. It was also recognized that in
the examples of the present invention, by performing continuous
rolling and controlling the temperature of the sheet bar in the
long direction and the width direction, the thickness precision and
the shape of the produced steel sheet are improved, and variations
in material properties are decreased.
[0542] Industrial Availability
[0543] According to the present invention, a cold rolled steel
sheet can be obtained, in which TS is greatly increased by press
forming and heat treatment while maintaining excellent deep
drawability in press forming. The cold-rolled steel sheet has the
excellent effect of industrially producing coated steel sheets by
electro-galvanization, hot-dip galvanization, alloying hot-dip
galvanization.
1 TABLE 1 Chemical Composition (mass %) Equation Equation Steel C N
Si Mn B Al Nb P S (1)'* (2)'** Remarks A 0.0009 0.011 0.01 0.12
0.0009 0.010 0.016 0.009 0.005 0.0023 -0.0001 Suitable example B
0.0020 0.015 0.01 0.11 0.0011 0.012 0.035 0.015 0.006 0.0021
-0.0003 Suitable example C 0.0005 0.009 0.01 0.09 0.0005 0.009
0.010 0.011 0.005 0.0022 -0.0001 Suitable example D 0.0020 0.021
0.01 0.50 0.0015 0.020 0.035 0.030 0.004 0.0035 -0.0003 Suitable
example E 0.0003 0.010 0.50 0.12 0.0006 0.011 0.099 0.045 0.010
0.0022 -0.0003 Suitable example F 0.0011 0.030 0.80 0.80 0.0011
0.028 0.025 0.009 0.005 0.0103 -0.0005 Suitable example G 0.0011
0.018 0.70 0.12 0.0008 0.012 0.018 0.008 0.005 0.0080 -0.0001
Suitable example H 0.0005 0.020 0.35 0.11 0.0014 0.020 0.025 0.007
0.005 0.0041 -0.0011 Suitable example I 0.0098 0.002 0.01 0.12
0.0007 0.038 0.055 0.009 0.005 -0.0269 0.0027 Comparative example J
0.0022 0.001 0.50 0.12 0.0008 0.012 0.001 0.008 0.005 -0.0064
0.0021 Comparative example K 0.0260 0.003 0.02 0.25 0.0001 0.035
0.001 0.013 0.007 -0.0154 0.0259 Comparative example L 0.0027 0.011
0.01 0.12 0.0007 0.014 0.001 0.009 0.005 0.0027 0.0026 Comparative
example *Equation (1)': N % - (14/93 .multidot. Nb % + 14/27
.multidot. Al % + 14/11 .multidot. B %) (the suitable range of this
invention is 0.0015 or more) **Equation (2)': C % - (0.5 .multidot.
12/93 .multidot. Nb %) (the suitable range of this invention is 0
or less)
[0544]
2 TABLE 2 Hot rolling condition Cooling Cold rolling condition
condition Cold Recrystalli- Presence of Slab Finisher after rolling
zation alloying heating delivery finish Coiling reduction Heating
annealing hot-dip temperature temperature rolling temperature ratio
rate temperature galvani- No. Steel (.degree. C.) (.degree. C.) (s,
.degree. C./s) (.degree. C.) (%) (.degree. C./s) (.degree. C.)
zation Remarks 1 A 1150 920 0.39, 58 600 80 15 840 Yes Suitable
example 2 B 1100 900 0.32, 53 550 80 15 840 Yes Suitable example 3
C 1130 650* 0.25, 15 600 80 15 840 No Suitable example 4 D 1110 900
0.35, 59 400 80 15 840 No Suitable example 5 E 1160 920 0.28, 70
650 80 15 840 Yes Suitable example 6 F 1150 910 0.39, 65 550 80 15
840 Yes Suitable example 7 G 1150 900 0.36, 58 500 80 15 840 No
Suitable example 8 H 1000 680* 0.26, 20 500 80 15 840 No Suitable
example 9 A 1150 920 1.13, 10 600 80 25 840 Yes Suitable example 10
E 1160 920 0.67, 15 650 80 25 840 Yes Suitable example 11 I 1140
920 0.35, 57 650 80 15 840 Yes Comparative example 12 J 1000 900
0.68, 7 550 80 15 840 Yes Comparative example 13 K 1150 880 0.38,
61 500 80 15 840 No Comparative example 14 L 1000 920 1.02, 9 600
80 15 940 No Comparative example *The finisher delivery temperature
was lower than Ar.sub.3 transformation point. The cooling
conditions after finish rolling include the cooling start time(s)
and the cooling rate (.degree. C./s).
[0545]
3 TABLE 3 Charge in tensile strength after forming-heat treatment
T.S. .DELTA.T.S. T.S. .DELTA.T.S. Product characteristics after
after after after Tensile Crystal heat heat heat heat strength
grain treatment treatment treatment treatment T.S. r diameter
.DELTA.El at 120.degree. C. at 120.degree. C. at 170.degree. C. at
170.degree. C. No. Steel (MPa) value (.mu.m) (%) (MPa) (MPa) (MPa)
(MPa) Remarks 1 A 360 1.8 18 1.6 435 75 455 95 Suitable example 2 B
365 1.7 12 1.2 460 95 470 105 Suitable example 3 C 355 2.3 19 1.8
420 65 440 85 Suitable example 4 D 390 1.9 14 1.3 480 90 490 100
Suitable example 5 E 430 1.8 18 1.7 520 90 530 100 Suitable example
6 F 460 1.6 18 1.7 575 115 595 135 Suitable example 7 G 430 1.8 13
1.4 540 110 555 125 Suitable example 8 H 390 2.1 19 1.7 490 100 510
120 Suitable example 9 A 352 1.9 22 2.4 414 62 440 88 Suitable
example 10 E 417 1.9 23 2.7 489 72 503 86 Suitable example 11 I 370
1.9 18 3.2 420 50 450 80 Comparative example 12 J 400 1.6 22 2.4
435 35 460 60 Comparative example 13 K 360 1.1 26 3.1 390 30 415 55
Comparative example 14 L 360 1.2 25 4.2 414 54 445 85 Comparative
example
[0546]
4 TABLE 4 Heat treatment temperature (.degree. C.) 120 200 300
Retention 2 82 114 133 time 5 86 119 136 (min.) 10 91 122 138 20 95
125 140
[0547]
5 TABLE 5 TS .times. r value .DELTA.TS Al % N/Al MPa MPa 0.020 0.75
775 58 0.036 0.42 762 55 0.049 0.31 753 42 0.072 0.21 720 25 0.080
0.19 719 19
[0548]
6TABLE 6 C Si Mn P S N Al Nb Ti V B N/Al + Nb + Ar.sub.3 Steel % %
% % % % % % % % % Ti + V + B)* .degree. C. A 0.0013 0.01 0.15 0.009
0.005 0.0135 0.011 -- -- -- -- 1.23 884 B 0.0012 0.50 0.50 0.011
0.004 0.0175 0.010 0.010 -- -- 0.0010 0.83 888 C 0.0008 0.01 0.01
0.009 0.003 0.0190 0.009 -- 0.015 0.015 0.0010 0.48 896 D 0.0025
0.01 0.09 0.005 0.003 0.0160 0.010 0.010 0.015 -- 0.0006 0.45 876 E
0.0016 0.55 0.80 0.04 0.003 0.0195 0.015 0.009 -- 0.018 0.0012 0.45
905 F 0.0011 0.75 0.75 0.02 0.002 0.0135 0.008 0.015 -- -- 0.0010
0.56 889 G 0.0032 0.01 0.14 0.009 0.003 0.0130 0.045 -- -- -- --
0.29 894 H 0.0019 0.45 0.45 0.04 0.005 0.0140 0.015 0.035 -- --
0.0012 0.27 882 *The concentrations of Ti, V and B not added are
shown by "--", and handled as "0" in calculation of N/(Al + Nb + Ti
+ V + B).
[0549]
7 TABLE 7 Hot rolling Hot-rolled Cold Finish rolling sheet rolling
Rough Presence Reduc- annealing Reduc- rolling of tion Coiling
Process tion SRT RDT FET lubrica- FDT ratio CT Temp. Batch ratio
No. Steel .degree. C. .degree. C. .degree. C. tion .degree. C.
.degree. C. % .degree. C. Continuous % 1 A 1150 900 810 Yes 660 95
510 790 Batch 80.0 2 B 1150 910 830 Yes 660 95 510 840 Continuous
80.0 3 C 1140 920 840 Yes 680 95 520 800 Batch 82.5 4 D 1180 900
820 Yes 650 95 500 830 Continuous 82.5 5 E 1190 930 850 Yes 690 95
530 810 Batch 80.0 6 F 1160 920 840 Yes 670 95 530 810 Batch 82.5 7
G 1120 910 820 Yes 670 95 520 800 Batch 80.0 8 H 1100 900 810 Yes
650 95 500 790 Batch 80.0 9 A 1080 900 810 Yes 660 95 500 790 Batch
82.5 10 B 1140 910 820 Yes 670 65 520 790 Batch 82.5 11 C 1250 1150
1070 Yes 910 95 750 850 Continuous 80.0 12 D 1170 900 810 Yes 660
95 510 No processing 80.0 Cold-rolled sheet Cooling: recrystalli-
cool Alloying zation rate Plating Reten- annealing before Sheet
Bath tion Temp. Time plating temp. temp. Temp. time No. Steel
.degree. C. s .degree. C./s .degree. C. .degree. C. .degree. C. s
Remarks 1 A 860 40 15 470 465 -- -- II 2 B 860 40 -- -- -- -- -- I
3 C 870 40 30 465 460 460 25 III 4 D 850 40 35 470 465 -- -- II 5 E
880 40 25 465 460 460 23 III 6 F 880 40 -- -- -- -- -- I 7 G 870 40
50 480 470 -- -- II 8 H 860 40 15 465 470 470 20 III 9 A 860 40 15
470 465 -- -- II 10 B 860 40 1 470 460 460 20 III 11 C 870 40 30
465 460 460 25 III 12 D 850 40 -- -- -- -- -- I SRT = slab heating
temperature, RDT = rough rolling delivery temperature, FET =
finisher entrance temperature, FDT = finisher delivery temperature,
CT = coiling temperature, I = cold-rolled steel sheet, II = hot-dip
galvanized steel sheet, III = galvanized steel sheet, III =
alloying hot-dip galvanized steel sheet
[0550]
8 TABLE 8 Tensile property Strain age before strain aging hardening
Dissolved TSxr property N YS TS El r value BH .DELTA.TS No. Steel %
MPa MPa % value MPa MPa MPa Remarks 1 A 0.0069 225 321 53 2.4 770
122 75 Example 2 B 0.0089 274 391 43 2.3 899 183 93 Example 3 C
0.0054 221 316 54 2.8 885 97 72 Example 4 D 0.0049 221 316 54 2.8
885 90 66 Example 5 E 0.0050 304 435 39 2.0 870 80 63 Example 6 F
0.0088 304 434 39 2.1 911 133 87 Example 7 G 0.0000 224 320 53 2.8
896 3 0 Comparative example 8 H 0.0000 284 405 42 2.1 851 2 0
Comparative example 9 A 0.0070 215 311 50 2.3 715 152 83
Comparative example 10 B 0.0082 274 391 43 1.9 743 143 94
Comparative example 11 C 0.0035 236 331 51 2.0 662 93 48
Comparative example 12 D 0.0041 241 336 51 2.1 706 85 49
Comparative example
[0551]
9 TABLE 9 Chemical Composition (% by mass) Steel C Si Mn P S Al N
N/Al Nb 12/93 Nb B Others A 0.0025 0.15 0.85 0.050 0.002 0.005
0.0120 2.40 0.0040 0.0005 0.0015 -- B 0.0050 0.51 1.20 0.004 0.001
0.008 0.0150 1.88 0.0070 0.0009 0.0009 -- C 0.0024 0.15 0.35 0.040
0.001 0.004 0.0120 3.00 0.0050 0.0006 0.0015 -- D 0.0150 0.15 0.88
0.010 0.002 0.010 0.0100 1.00 0.0102 0.0013 0.0012 -- E 0.0023 0.01
1.40 0.005 0.002 0.011 0.0120 1.09 0.0070 0.0009 0.0010 -- F 0.0023
0.05 1.35 0.045 0.001 0.007 0.0120 1.71 0.0080 0.0010 0.0020 Mo:
0.15 G 0.0025 0.15 1.25 0.007 0.001 0.004 0.0110 2.75 0.0070 0.0009
0.0015 Ti: 0.013 H 0.0025 0.25 1.25 0.008 0.001 0.011 0.0140 1.27
0.0060 0.0008 0.0014 Cu: 0.50, Ni: 0.020 I 0.0024 0.15 1.25 0.005
0.003 0.011 0.0150 1.36 0.0089 0.0010 0.0018 Ni: 0.05 V: 0.02 J
0.022 0.15 1.21 0.008 0.002 0.005 0.0120 2.40 0.0076 0.0011 0.0020
Cu: 0.10 Ni: 0.05 K 0.070 0.25 1.50 0.015 0.003 0.055 0.0040 0.07
0.0090 0.0012 0.0021 -- L 0.0025 0.15 0.85 0.050 0.002 0.005 0.0120
2.40 0.0001 0.0000 <0.0001 -- M 0.0025 0.15 0.85 0.050 0.002
0.005 0.0120 2.40 0.0050 0.0006 <0.0001 -- O 0.0025 0.17 0.81
0.008 0.001 0.007 0.0120 1.71 0.110 0.0142 0.0015 -- P 0.0026 0.17
0.82 0.007 0.001 0.007 0.0400 5.71 0.0070 0.0009 0.0017 -- Q 0.0025
0.18 0.81 0.008 0.001 0.007 0.0120 1.71 0.0080 0.0010 0.0022 Ca:
0.0035
[0552]
10 TABLE 10 Hot rolling step Rough +TC,19 Cold rolling step Slab
rolling Finishing rolling Cold Thickness heating Thickness Bonding
Ar.sub.3 Finisher Thickness Coiling rolling of Steel temp. of of
transformation temperature of temperature reduction cold-rolled
sheet SRT sheet bar sheet point FDT hot-rolled sheet CT ratio sheet
No. Steel .degree. C. mm bars .degree. C. .degree. C. mm .degree.
C. % mm 1 A 1210 28 No 873 900 4.0 550 83 0.70 2 B 1230 28 No 837
900 4.0 530 81 0.75 3 C 1220 28 No 908 910 4.0 550 83 0.70 4 D 1180
28 No 851 890 4.0 530 83 0.70 5 E 1190 28 No 801 890 4.0 520 78
0.70 6 F 1180 25 No 827 890 4.5 500 80 1.00 7 G 1180 25 No 820 890
6.0 570 80 1.20 8 H 1180 25 No 824 890 4.0 610 84 0.80 9 I 1190 28
No 819 890 4.0 550 83 0.65 10 J 1210 28 No 823 870 4.0 450 83 0.70
11 K 1200 28 No 808 880 4.0 520 83 0.70 12 L 1200 28 No 873 900 4.0
550 83 0.70 13 M 1200 25 No 873 900 4.0 550 83 0.70 14 O 1210 25 No
856 890 4.0 520 83 0.70 15 P 1240 25 No 854 880 4.0 520 83 0.70 16
Q 1220 25 No 856 900 4.0 530 83 0.70 17 A 950 25 No 870 700 4.0 540
83 0.70 18 A 1200 28 No 873 900 4.0 780 83 0.70 19 A 1190 28 No 873
905 4.0 540 83 0.70 20 A 1190 28 No 873 890 4.0 550 83 0.70
Temperature Cold-rolled sheet annealing step region Continuous
annealing Steel of Annealing Retention Cooling Cooling stop Temper
rolling sheet .alpha.-.gamma. coexistence temp. time rate rate
Elongation No. Steel .degree. C. .degree. C. s .degree. C./s
.degree. C. Type % Remarks 1 A 860 to 950 880 35 30 350 Skin pass
0.5 Example of this invention 2 B 835 to 935 870 30 35 350 Skin
pass 0.5 Example of this invention 3 C 865 to 955 885 25 25 350
Skin pass 0.5 Example of this invention 4 D 845 to 922 885 40 45
350 -- -- Example of this invention 5 E 825 to 906 875 25 40 450
Leveler 0.5 Example of this invention 6 F 855 to 930 885 30 35 450
Skin pass + Leveler 0.5 Example of this invention 7 G 835 to 918
875 45 30 450 Skin pass + Leveler 0.7 Example of this invention 8 H
835 to 923 880 25 45 450 Skin pass + Leveler 0.7 Example of this
invention 9 I 830 to 915 865 25 30 300 Skin pass + Leveler 0.7
Example of this invention 10 J 830 to 910 870 25 30 300 Skin pass +
Leveler 0.7 Example of this invention 11 K 755 to 895 870 25 30 300
Skin pass + Leveler 0.5 Comparative example 12 L 860 to 950 880 35
30 300 Skin pass + Leveler 0.5 Comparative example 13 M 860 to 950
880 35 30 300 Skin pass + Leveler 0.5 Comparative example 14 O 850
to 940 890 30 35 300 Skin pass + Leveler 0.5 Comparative example 15
P 850 to 940 890 30 35 300 Skin pass 0.7 Comparative example 16 Q
850 to 940 880 30 35 300 Skin pass 0.7 Example of this invention 17
A 855 to 950 880 30 30 300 Skin pass 0.7 Comparative example 18 A
865 to 950 870 35 35 300 Skin pass 0.7 Comparative example 19 A 860
to 950 820 25 40 300 Skin pass 0.7 Comparative example 20 A 860 to
950 870 30 5 300 Skin pass 0.7 Comparative example
[0553]
11 TABLE 11 Amount Steel sheet Property of of Second product sheet
Strain age hardening dissolved Ferrite phase Tensile property
property Aging Steel N Area Grain AF Mean Amount of resistance
sheet % ratio diameter area ratio YS TS El r BH .DELTA.TS .DELTA.Y
- El No. Steel by mass % .mu.m % MPa MPa % value MPa MPa % Remarks
1 A 0.0090 92 1 8 315 448 38 1.8 85 75 0.0 Example of this
invention 2 B 0.0100 95 8 5 355 510 34 1.7 90 80 0.0 Example of
this invention 3 C 0.0085 92 7 8 295 420 41 1.8 80 75 0.0 Example
of this invention 4 D 0.0095 92 6 8 334 475 36 1.4 95 80 0.0
Example of this invention** 5 E 0.0100 93 7 7 325 465 37 1.7 88 75
0.0 Example of this invention 6 F 0.0105 90 7 10 385 550 35 1.5 102
75 0.0 Example of this invention 7 G 0.0105 88 6 12 315 455 39 1.8
88 72 0.0 Example of this invention 8 H 0.0095 92 10 8 325 460 37
1.8 98 75 0.0 Example of this invention 9 I 0.0090 90 5 10 320 465
37 1.8 97 85 0.0 Example of this invention 10 J 0.0105 88 5 12 420
595 29 1.4 105 85 0.0 Example of this invention** 11 K 0.0002 75 7
25 370 545 25 1.1 45 34 1.5 Comparative example 12 L 0.0080 98 15 2
290 420 35 1.3 15 10 1.5 Comparative example 13 M 0.0085 97 8 3 290
410 33 1.2 15 7 1.0 Comparative example 14 O 0.0003 98 6 2 320 455
33 1.1 5 15 0.8 Comparative example 15 P 0.0190 95 7 5 310 440 37
1.6 90 70 2.8 Comparative example 16 Q 0.0098 92 7 8 315 450 39 1.7
90 70 0 Example of this invention 17 A 0.0009 98 22 2 275 420 33
1.1 0 15 0.5 Comparative example 18 0.0040 95 25 5 275 410 34 1.3
30 10 0.5 Comparative example 19 0.0050 100 12 0 270 380 33 1.3 25
12 0.8 Comparative example 20 0.0030 99 22 1 265 385 31 1.3 20 10
0.9 Comparative example AF: Acicular ferrite **Hot-dip
galvanization
[0554]
12 TABLE 12 Chemical Composition (% by mass) Steel C Si Mn P S Al N
N/Al Nb 12/93 Nb B Others N 0.0085 0.005 0.55 0.009 0.005 0.015
0.0126 0.84 0.008 0.0010 0.0015 --
[0555]
13 TABLE 13 Rough rolling Finish rolling Coiling Cold rolling step
Slab Thick- Ar.sub.3 Finisher Thickness Coiling Cold Thickness
heat- ness Bond- trans- temp- of temp- rolling of ing of ing for-
era- hot- era- reduc- cold- Steel temp. sheet of mation ture rolled
ture tion rolled sheet SRT bar sheet point FDT sheet CT ratio sheet
No. Steel .degree. C. mm bars .degree. C. .degree. C. mm .degree.
C. % mm 2-1 N 1250 25 Yes* 870 890 4.5 570 69 1.4 2-2 " 1230 25
Yes* 870 890 4.5 570 69 1.4 2-3 " 1240 25 Yes** 870 890 4.5 570 69
1.4 Temperature Cold-rolled sheet annealing step region Continuous
annealing .alpha.-.gamma. Anneal- Cooling Steel of co- ing Re-
Cooling stop Temper Rolling sheet existence temp. tention rate
temp. Elongation No. Steel .degree. C. .degree. C. time s .degree.
C./s .degree. C. Type % Remarks 2-1 N 840 to 900 880 30 45 300 Skin
0.5 Example pass of this inven- tion 2-2 " 840 to 900 880 35 45 310
Skin 0.5 Example pass of this inven- tion 2-3 " 840 to 900 880 30
45 310 Skin 0.5 Example pass of this inven- tion *Lubrication
rolling **Lubrication rolling, using a sheet bar heater and an edge
heater
[0556]
14 TABLE 14 Property of product Property sheet Strain Steel sheet
of product Variation age Amount Second sheet in hardening of phase
Tensile tensile property Aging dis- Ferrite AF property property *
Amount re- Steel solved Area Grain area Mean of sistance sheet N %
ratio diameter ratio YS TS El r .delta.YS .delta.TS .delta.El BH
.DELTA.TS .DELTA.Y-El No. Steel by mass % .mu.m % MPa MPa % value
MPa MPa % MPa MPa % Remarks 2-1 N 0.0120 91 7 9 315 445 37 1.6 10
10 5 15 15 0 Example of this invention 2-2 N 0.0120 91 7 9 319 447
37 1.6 5 5 2 7 6 0 Example of this invention 2-3 N 0.0125 92 7 8
318 450 38 1.5 3 5 1 8 5 0 Example of this invention * .delta.YS,
.delta.TS, .delta.El: = (Maximum - Minimum) ** AF: Acicular
ferrite, M: Martensite, B: Bainite, P: Pearlite
[0557]
15 TABLE 15 Transformation Chemical Composition (% by mass) point
(.degree. C.) Steel C Si Mn P S Al N N/Al Other Ac.sub.1 Ac.sub.3 A
0.050 0.01 0.55 0.040 0.001 0.004 0.0120 3.00 -- 735 858 B 0.080
0.02 0.30 0.060 0.001 0.008 0.0153 1.91 -- 749 866 C 0.070 0.15
0.45 0.040 0.001 0.002 0.0118 5.90 REM: 0.0015 740 862 D 0.055 0.15
0.88 0.050 0.002 0.006 0.0095 1.58 Ca: 0.0020 739 863 E 0.030 0.01
1.40 0.005 0.002 0.011 0.0120 1.09 Ca: 0.0025, 706 832 REM: 0.0020
F 0.050 0.05 0.55 0.045 0.001 0.007 0.0123 1.76 Mo: 0.02 739 862 G
0.065 0.02 0.75 0.045 0.001 0.004 0.0114 2.85 Ti: 0.013, 735 855 B:
0.0005 H 0.050 0.01 0.85 0.008 0.001 0.008 0.0138 1.73 Cu: 0.50,
717 834 Ni: 0.20, Cr: 0.20 I 0.040 0.01 0.55 0.005 0.003 0.011
0.0151 1.37 Ni: 0.05, 717 842 V: 0.02 J 0.120 0.15 0.95 0.008 0.002
0.005 0.0118 2.36 Cu: 0.10, 719 830 Ni: 0.05, Nb: 0.007 K 0.010
0.25 0.96 0.015 0.003 0.002 0.0042 2.10 -- 723 857 L 0.080 0.15
0.85 0.050 0.002 0.045 0.0040 0.09 -- 739 859
[0558]
16 TABLE 16 Hot rolling step Rough rolling Finish rolling Cooling
Coiling Cold rolling step Slab Thick- Finisher Thickness after
rolling Coiling Cold Thickness heat- ness Bond- temp- of Start-
Cool- temp- rolling of ing of ing era- hot- ing ing era- reduc-
cold- Steel temp. sheet of ture rolled time rate ture tion rolled
sheet SRT bar sheet FDT sheet .DELTA.t V CT ratio sheet No. Steel
.degree. C. mm bars .degree. C. mm s .degree. C./s .degree. C. % mm
1 A 1210 30 No 850 4.0 0.2 50 550 83 0.70 2 B 1230 30 No 870 4.0
0.2 50 530 81 0.75 3 C 1220 30 No 840 4.0 0.1 50 550 83 0.70 4 D
1180 30 No 850 4.0 0.5 50 530 83 0.70 5 E 1190 30 No 850 4.0 0.2 45
520 83 0.70 6 F 1180** 35 No 864 4.5 0.2 45 500 78 1.00 7 G 1180 35
No 860 6.0 1.2 45 570 80 1.20 8 H 1180 35 No 830 4.0 0.3 45 610 80
0.80 9 I 1190 30 Yes 840 4.0* 0.3 45 550 84 0.65 10 J 1210 30 No
840 4.0 0.3 50 450 83 0.70 11 K 1200 30 No 850 4.0 0.3 50 520 83
0.70 12 L 1200 30 No 880 4.0 0.2 50 550 83 0.70 13 A 950 35 No 750
4.0 0.5 50 540 83 0.70 14 A 1200 35 No 880 4.0 0.3 50 780 83 0.70
15 A 1190 35 No 870 4.0 0.2 50 540 83 0.70 16 A 1200 35 No 880 4.0
0.2 50 550 83 0.70 17 B 1190 30 No 840 4.0 0.3 50 520 83 0.70 18 G
1180 30 No 850 4.0 0.3 50 520 83 0.70 Cold-rolled sheet annealing
step Over aging Box Continuous annealing Residence time annealing
Cooling in temperature Temper Steel Annealing Annealing Cooling
stop region of rolling sheet temp. temp. Retention rate temp.
350.degree. C. or Elongation No. Steel .degree. C. .degree. C. time
s .degree. C./s .degree. C. higher ** s % 1 A 700 790 40 45 270 --
0.2 2 B 680 780 30 50 270 -- 0.2 3 C 700 800 30 55 270 -- 0.2 4 D
690 800 45 45 250 -- 0.2 5 E 720 770 25 45 200 -- 0.5 6 F 700 780
35 55 200 -- -- 7 G 700 830 50 60 300 -- 0.5 8 H 700 810 25 50 400
50 0.5 9 I 700 790 25 45 250 -- 0.5 10 J 700 770 25 30 250 -- 0.5
11 K 700 800 25 30 250 -- 0.5 12 L 700 790 35 30 270 -- 0.7 13 A
700 800 30 30 270 -- 0.7 14 A 700 830 35 35 250 -- 0.7 15 A 700 680
25 40 250 -- 0.5 16 A -- 820 29 35 250 -- 0.5 17 B 710 780 35 70
450 -- **** 18 G 720 780 40 60 450 -- **** *Lubrication rolling
**Residence time at 350 .degree. C. to the cooling stop temperature
***Charging of hot slab ****Temper rolling with an elongation of
0.5% after hot-dip galvanization
[0559]
17 TABLE 17 Steel Sheet Strain age Amount Second Property of
hardening of phase product sheet property Aging dis- Ferrite AF
Tensile property Amount re- Steel solved Area Grain area Other Mean
of sistance sheet N % ratio diameter ratio phase: YS TS El r BH
.DELTA.TS .DELTA.Y-El No. Steel by mass % .mu.m % type MPa MPa %
value MPa MPa % Remarks 1 A 0.0075 86 7 12 B 290 550 35 1.6 80 80
0.0 Example of this invention 2 B 0.0055 92 8 7 P 295 556 35 1.5 90
85 0.0 Example of this invention 3 C 0.0055 90 7 10 -- 285 555 34
1.6 85 75 0.0 Example of this invention 4 D 0.0075 89 6 10 B 295
565 34 1.6 95 80 0.0 Example of this invention 5 E 0.0110 91 7 9 --
315 605 33 1.7 85 75 0.0 Example of this invention 6 F 0.0110 90 7
10 -- 309 595 35 1.6 100 75 0.0 Example of this invention 7 G
0.0095 85 6 14 B 314 605 35 1.6 85 70 0.0 Example of this invention
8 H 0.0078 87 10 12 B 318 615 33 1.6 96 75 0.0 Example of this
invention 9 I 0.0086 86 5 12 P 276 525 37 1.5 90 80 0.0 Example of
this invention 10 J 0.0095 84 5 15 P 324 620 33 1.5 102 90 0.0
Example of this invention 11 K 0.0003 98 7 1 P 235 3550 38 1.4 45
34 1.5 Comparative example 12 L 0.0005 93 15 7 -- 241 370 39 1.3 15
10 1.5 Comparative example 13 A 0.0009 98 22 0.5 P 285 430 32 1.1 0
15 0.5 Comparative example 14 " 0.0009 95 25 5 -- 285 445 32 1.2 30
10 0.5 Comparative example 15 " 0.0010 100 12 0 -- 315 480 34 1.4
25 12 0.8 Comparative example 16 " 0.0020 99 22 1 -- 350 545 25 1.0
20 10 0.9 Comparative example 17 B 0.0060 92 8 7 P 297 555 35 1.5
92 85 0 Example of this invention 18 G 0.0055 92 8 7 P 296 557 34
1.5 90 85 0 Example of this invention M: Martensite, B: Bainite, P:
Pearlite
[0560]
18 TABLE 18 Transformation Chemical Composition (% by mass) point
(.degree. C.) Steel C Si Mn P S Al N N/Al Other Ac.sub.1 Ac.sub.3 M
0.052 0.01 0.60 0.035 0.001 0.002 0.0125 6.25 -- 735 855
[0561]
19 TABLE 19 Rough Cooling rolling Finish rolling after Coiling Cold
rolling step Slab Thick- Finisher Thickness rolling Coiling Cold
Thickness heat- ness Bond- temp- of Cool temp- rolling of ing of
ing era- hot- Start- ing era- reduc- cold- Steel temp. sheet of
ture rolled ing rate ture tion rolled sheet SRT bar sheet FDT sheet
time V CT ratio sheet No. Steel .degree. C. mm bars .degree. C. mm
.DELTA.t s .degree. C./s .degree. C. % mm 2-1 M 1180 30 No 850 4.5
0.3 45 540 64 1.6 2-2 " 1200 30 Yes* 850 4.5 0.3 45 540 64 1.6 2-3
" 1190 30 Yes** 855 4.5 0.3 45 540 64 1.6 Over aging Residence Box
Continuous annealing time annealing Cooling in temperature Temper
Steel Annealing Annealing Re- Cooling stop region of rolling sheet
temp. temp. tention rate temp. 350.degree. C. or Elongation No.
Steel .degree. C. .degree. C. time s .degree. C./s .degree. C.
higher ** s % Remarks 2-1 M 700 790 40 45 270 -- 0.2 Example of
this invention 2-2 " 700 790 40 45 270 -- 0.2 Example of this
invention 2-3 " 700 790 40 45 270 -- 0.2 Example of this invention
*Lubrication rolling **Using a sheet bar heater, an edge heater
[0562]
20 TABLE 20 Steel sheet Property of Amount Second phase product
sheet of Ferrite AF Tensile property Steel dissolved Area Grain
area Other Mean sheet N % ratio diameter ratio phases: YS TS El r
No. Steel by mass % .mu.m % type ** MPa MPa % value 2-1 M 0.0075 85
7 12 B 295 551 35 1.6 2-2 M 0.0075 85 7 12 B 295 552 35 1.6 2-3 M
0.0075 85 7 12 B 292 553 35 1.6 Property of Strain age product
sheet hardening property Variation in Amount Aging Steel tensile
property * of resistance sheet .delta.YS .delta.TS .delta.El BH
.DELTA.TS .DELTA.Y-El No. MPa MPa % MPa MPa % Remarks 2-1 25 15 2
80 80 0 Example of this invention 2-2 20 10 1 80 80 0 Example of
this invention 2-3 15 8 1 80 80 0 Example of this invention *
.delta.YS, .delta.TS, .delta.El: = (Maximum - Minimum) ** M:
Martensite, B: Bainite, P: Pearlite
* * * * *