U.S. patent application number 13/805144 was filed with the patent office on 2013-04-11 for ultra high strength cold rolled steel sheet having excellent ductility and delayed fracture resistance and method for manufacturing the same.
This patent application is currently assigned to JFE STEEL CORPORATION. The applicant listed for this patent is Kohei Hasegawa, Masataka Yoshino. Invention is credited to Kohei Hasegawa, Masataka Yoshino.
Application Number | 20130087257 13/805144 |
Document ID | / |
Family ID | 45402222 |
Filed Date | 2013-04-11 |
United States Patent
Application |
20130087257 |
Kind Code |
A1 |
Yoshino; Masataka ; et
al. |
April 11, 2013 |
ULTRA HIGH STRENGTH COLD ROLLED STEEL SHEET HAVING EXCELLENT
DUCTILITY AND DELAYED FRACTURE RESISTANCE AND METHOD FOR
MANUFACTURING THE SAME
Abstract
An ultra.-high-strength cold-rolled steel sheet with excellent
ductility and delayed fracture resistance includes 0.15% to 0.75 C.
1.0% to 3.0% Si, 1.5% to 2.5% Mn, 0.05% or less P, 0.02% or less 5,
0.01% to 0.05% Al, and less than 0.005% N on a mass ratio, the
remainder being Fe and =avoidable impurities, the
ultra-high-strength cold-rolled steel sheet having a metal
microstructure including 40% to 85% of a tempered martensite phase
and 15% to 60% of a ferrite phase on a volume fraction basis and a
tensile strength of 1320 Mtn or more.
Inventors: |
Yoshino; Masataka; (Tokyo,
JP) ; Hasegawa; Kohei; (Tokyo, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
Yoshino; Masataka
Hasegawa; Kohei |
Tokyo
Tokyo |
|
JP
JP |
|
|
Assignee: |
JFE STEEL CORPORATION
Tokyo
JP
|
Family ID: |
45402222 |
Appl. No.: |
13/805144 |
Filed: |
June 24, 2011 |
PCT Filed: |
June 24, 2011 |
PCT NO: |
PCT/JP2011/065135 |
371 Date: |
December 18, 2012 |
Current U.S.
Class: |
148/648 ;
148/320; 148/330 |
Current CPC
Class: |
C21D 8/005 20130101;
C21D 2211/005 20130101; C21D 9/46 20130101; C22C 38/14 20130101;
C21D 2211/008 20130101; C22C 38/04 20130101; C21D 8/0436 20130101;
C22C 38/001 20130101; C21D 8/0447 20130101; C21D 8/00 20130101;
C22C 38/06 20130101; C22C 38/02 20130101; C21D 6/00 20130101; C22C
38/12 20130101 |
Class at
Publication: |
148/648 ;
148/320; 148/330 |
International
Class: |
C21D 8/00 20060101
C21D008/00; C22C 38/12 20060101 C22C038/12; C22C 38/00 20060101
C22C038/00; C22C 38/04 20060101 C22C038/04; C22C 38/02 20060101
C22C038/02; C22C 38/14 20060101 C22C038/14; C22C 38/06 20060101
C22C038/06 |
Foreign Application Data
Date |
Code |
Application Number |
Jun 30, 2010 |
JP |
2010-148531 |
Claims
1. An ultra-high-strength cold-rolled steel sheet with excellent
ductility and delayed fracture resistance, comprising 0.15% to
0.25% C, 1.0% to 3.0% Si, 1.5% to 2.5% Mn, 0.05% or less P, 0.02%
or less S, 0.01% to 0.05% Al, and less than 0.005% N on a mass
ratio, the remainder being Fe and unavoidable impurities, the
ultra-high-strength cold-rolled steel sheet having a microstructure
comprising 40% to 85% of a tempered martensite phase and 15% to 60%
of a ferrite phase on a volume fraction basis and a tensile
strength of 1320 MPa or more.
2. The cold-rolled steel sheet according to claim 1, further
comprising one or more of 0.1% or less Nb, 0.1% or less Ti, and 5
ppm to 30 ppm B on a mass ratio.
3. The cold-rolled steel sheet according to claim 1, having a total
elongation of 12% or more.
4. A method for manufacturing an ulta-high-strengh cold-rolled
steel sheet having excellent ductility and delayed fracture
resistance, comprising; heating a steel slab having the chemical
composition specified in claim 1 to 1200.degree. C. or higher;
hot-rolling the steel slab at a finish rolling end temperature of
800.degree. C. or higher; pickling the steel; cold-rolling the
steel; continuously annealing the steel such that the steel is held
at a temperature ranging from the Ac.sub.1 transformation
temperature to Ac.sub.3 transformation temperature thereof for 30 a
to 1200 s, cooled to a temperature of 600.degree. C. to 800.degree.
C. at an average cooling rate of 100.degree. C./s or less, and then
cooled to 100.degree. C. or lower at an average cooling rate of
100.degree. C./s to 1000.degree. C./s; and tempering the steel such
that the steel is reheated and held at a temperature of 100.degree.
C. to 300.degree. C. for 120 s to 1800 s.
5. The cold-rolled steel sheet according to claim 2, having a total
elongation of 12% or more.
6. A method for manufacturing an ultra-high-strength cold-rolled
steel sheet having excellent ductility and delayed fracture
resistance, comprising; heating a steel slab having the chemical
composition specified in claims 2 to 1200.degree. C. or higher;
hot-rolling the steel slab at a finish roiling end temperature
of800.degree. C. or higher; pickling the steel; cold-rolling the
steel; continuously annealing the steel such that the steel is held
at a temperature ranging from the Ac1 transformation temperature to
Ac3 transformation temperature thereof for 30 s to 1200 s, cooled
to a temperature of 600.degree. C. to 800'C at an average cooling
rate of 100.degree. Cis or less, and then cooled to 100.degree. C.
or lower at an average cooling rate of 100.degree. C/.s to
1000.degree. C./s; and tempering the steel such that the steel is
reheated and held at a temperature of 100.degree. C. to 300.degree.
C. for 120 s to 1800 a.
Description
RELATED APPLICATIONS
[0001] This is a .sctn.371 of International Application No.
PCT/JP2011 /065135, with an international filing date of Jun. 24,
2011 (WO 2012/002520 Al, published Jan. 5, 2012), which is based on
Japanese Patent Application No. 2010-148531, filed Jun. 30, 2010,
the subject matter of which is incorporated herein by
reference.
TECHNICAL FIELD
[0002] This disclosure relates to an ultra-high-strength
cold-rolled steel sheet which has an excellent strength-ductility
balance and excellent delayed fracture resistance and which is a
material suitable for use principally in ultra-high-strength
automobile structural parts such as center pillars and door impact
beams for automobiles and also relates to a method for
manufacturing the same.
BACKGROUND
[0003] In recent years, in Europe, regulations on CO.sub.2
emissions from automobiles, which are mobile CO.sub.2 emission
sources, have been tightened because of concerns about global
warming due to increasing CO.sub.2 emissions and therefore the
improvement of automobile fuel efficiency has been strongly
required. The lightening of car bodies is effective in improving
fuel efficiency. However, since the safety of occupants needs to be
ensured, the crash safety of light-weight car bodies needs to be
more improved than ever. To cope with two requirements, that is,
the lightening of car bodies and the ensuring of crash safety, the
gauge reduction of steel sheets used is being attempted using
materials with high specific strength. In recent years,
high-strength steel sheets with a tensile strength of 980 to 1180
MPa have been actively used for auto-mobile structural parts such
as center pillars and door impact beams. However, there are
increasing demands for lightweight car bodies and therefore
attempts are being made to manufacture more lightweight car bodies
using steel sheets stronger than 1180 MPa class steel sheets,
[0004] Since automobile structural parts are usually manufactured
by press molding, the ductility of materials significantly affects
the press formability thereof in view of automobile crash safety,
residual ductility after press molding is important. Since the
ductility of steel sheets usually decreases with an increase in
strength, press formability and residual ductility after press
molding decrease with an increase in strength. In high-strength
materials with a tensile strength of greater than 980 MPa, there
are concerns about delayed fracture due to residual ductility after
press molding and hydrogen coming from surroundings. Therefore, to
use high-strength cold-rolled steel sheets for the above automobile
structural parts, the high-strength cold-rolled steel sheets need
to have high press formability, high ductility, and excellent
delayed fracture resistance
[0005] Various proposals have been made to cope with these
requirements.
[0006] For example, Japanese Unexamined Patent Application
Publication No. 2005-163055 discloses an example in which a steel
sheet assumed to have a tensile strength of 1350 MPa and a tempered
martensite single-phase microstructure is obtained by quench and
tempering although the percentages of phases are not described
therein. However, the total elongation of the steel sheet is small,
7%. Therefore, it is extremely difficult to manufacture automobile
safety parts from the steel sheet by pressing. The martensite
single-phase microstructure is probably obtained by quenching and
therefore the steel sheet probably has a seriously had shape. This
case needs a step of correcting the shape thereof after annealing
and therefore is not preferable in terms of manufacture.
[0007] Japanese Unexamined Patent Application Publication No.
2006-307325 discloses a TRIP (Transformation-induced Plasticity)
steel sheet which has high strength and ductility and which is
obtained by making use of strain-induced transformation, that is,
the transformation of retained austenite into martensite by strain
during deformation. That steel sheet contains 0.3% to 2% Al on a
mass basis to ensure the amount of retained austenite necessary to
develop a TRIP effect. A large amount of Al causes a problem that
casting defects are likely to he caused. To allow retained
austenite to remain in a microstructure, isothermal holding needs
to be performed at a temperature not lower than the its
transformation temperature in the course of cooling from the
annealing temperature, which results in an increased number of
manufacturing steps. Since the change in rate of cooling to the
temperature of isothermal holding during operation causes a
significant change in material quality, operating conditions needs
to be strictly controlled to stably manufacture steel sheets with a
certain level of quality, which is not preferable in terms of
manufacture.
[0008] It could therefore be helpful to provide an
ultra-high-strength cold-rolled steel sheet which has excellent
delayed fracture resistance and a tensile strength of 1320 MPa or
more and which does not excessively contain a transition metal
element, such as V or Mo, causing a significant increase in
alloying cost or Al, which may possibly cause casting defects, and
to provide a method for manufacturing the ultra-high-strength
cold-rolled steel sheet.
SUMMARY
[0009] To obtain a conventional ultra-high-strength cold-rolled
steel sheet with a tensile strength of 1320 MPa or more, a
microstructure needs to be transformed into a martensite
single-phase microstructure by quenching. In the case where a
microstructure is a martensite single-phase, sufficient ductility
cannot be achieved. Even it an attempt is made to increase the
ductility by tempering subsequent to quenching, the strength is
reduced and ductility is apt not to be increased so much because of
the recovery of a dislocation microstructure in a martensite phase
and the coarsening of a carbide such as Fe.sub.3C, precipitated in
the martensite phase.
[0010] On the other hand, to develop high ductility, many TRIP
steels have been invented by making use of file strain-induced
transformation of a retained austenite. However, to develop a TRIP
effect, a large amount of an alloying element needs to be used to
increase the stability of austenite and isothermal holding needs to
be precisely performed at temperature not lower than the Ms
transformation temperature in the course of cooling from the
annealing temperature, which is not preferable in terms of
manufacturing stability and manufacturing costs.
[0011] In view of delayed fracture resistance, hydrogen-trapping
sites, which cause delayed fracture, are preferably diminished as
much as possible. Martensite phases are preferably diminished as
much as possible because a large number of dislocations serving as
hydrogen-trapping sites are introduced into the martensite phases
during crystallographic transformation from austenitic phases.
Retained austenite, which contributes to an increase in ductility,
is known to serve as a hydrogen-trapping site like a dislocation
and is present on a grain boundary in the form of a film.
Therefore, the penetration of hydrogen into retained austenite may
possibly cause grain boundary fracture to reduce delayed fracture
resistance. Thus, it is not preferred that a metal microstructure
contains retained austenite,
[0012] We discovered that the balance between tensile strength and
ductility can be controlled such that a microstructure is converted
into a microstructure containing a tempered martensite phase and a
ferrite phase and the volume fraction of the tempered martensite
phase is varied. We discovered a technique in which a steel sheet
with ultra-high strength is obtained such that the hardness of the
tempered martensite phase and that of the ferrite phase are
increased by the addition of C and Si the volume fraction of an
untempered martensite phase is reduced. We found that an
ultra-high-strength steel sheet with high ductility can be
obtained.
[0013] In addition, we discovered that the density of dislocations
in a microstructure can be significantly reduced as compared with a
smart-ensile single-phase microstructure by precipitating a ferrite
phase containing substantially no dislocation in the microstructure
and the amount of hydrogen permeating through steel can be
significantly reduced by diminishing hydrogen-trapping sites. We
thus found that delayed fracture resistance can be increased.
[0014] Furthermore, we found that in view of manufacturing steps,
it is effective the annealing temperature and the course of cooling
are appropriately controlled during annealing and cooling
subsequent to cold rolling and tempering heat treatment is
performed at a temperature of 100.degree. C. to 300.degree. C.
[0015] We thus provide: [0016] (1) An ultra-high-strength
cold-rolled steel sheet with excellent ductility and delayed
fracture resistance contains 0.15% to 0.25% C. 1.0% to 3.0% Si,
1.5% to 2.5% Mn, 0.05% or less P. 0.02% or less S. 0.01% to 0.05%
Al, and less than 0.005% N on a mass ratio, the remainder being Fe
and unavoidable impurities, and has a metal microstructure
containing 40% to 85% of a tempered martensite phase and 15% to 60%
of a terrific phase on a volume fraction basis and a tensile
strength of 1320 MPa or more. [0017] (2) The ultra-high-strength
cold-roiled steel sheet with excellent ducti thy and delayed
fracture resistance specified in item (1) further contains one or
more of 0.1% or less Nb, 0.1% or less Ti, and 5 ppm to 30 ppm B on
a mass ratio, [0018] (3) The ultra-high-strength cold-rolled steel
sheet with excellent ductility and delayed fracture resistance
specified in Item (1) or (2) has a total elongation of 12% or more.
[0019] (4) A method for manufacturing an ultra-high-strength
cold-rolled steel sheet having excellent ductility and delayed
fracture resistance includes heating a steel slab having the
chemical composition specified in Item (1) or (2) to 1200.degree.
C. or higher; hot-rolling the steel slab at a finish rolling end
temperature of 800.degree. C. or higher; pickling the steel;
cold-rolling the steel; continuously annealing the steel in such a
manner that the steel is held at a temperature ranging from the Ac
.sub.1 transforrnation temperature to Ac.sub.3 transformation
temperature thereof for 30 s to 1200 s, is cooled to a temperature
of 600.degree. C. to 800.degree. C. at an average cooling rate of
100.degree. C./s or less, and is then cooled to 100.degree. C. or
lower at an average cooling rate of 100.degree. C./s to
1000.degree. C./s; and tempering the steel in such a manner that
the steel is reheated and is held at a temperature of 100.degree.
C. to 300.degree. C. for 120 s to 1800 s.
[0020] Our cold-rolled steel sheet has extremely high tensile
strength, high ductility, and therefore excellent workability.
Parts fortned from the cold-rolled steel sheet have resistance to
delayed fracture due to hydrogen coming from surroundings, that is
excellent delayed fracture resistance. For example, a tensile
strength of 1320 MPa or more, a total elongation of 12% or more,
and such delayed fracture resistance that fracture does not occur
for 100 hours in a 25.degree. C. hydrochloric acid environment with
a pH of 3 can be readily achieved, Furthermore, a cold-rolled steel
sheet having such excellent properties as described above can be
stably manufactured by our method.
[0021] The following sheet can be stably manufactured: an
ultra-high-strength cold-rolled steel sheet which has a tensile
strength of 1320 MPa or more and which exhibits excellent
workability during forming. Parts formed from the cold-rolled steel
sheet by press molding have resistance to delayed fracture due to
hydrogen coming from surroundings, that is, excellent delayed
fracture resistance. Ultra-high-strength parts, such as automobile
safety parts including center pillars and impact beams, resistant
to delayed fracture can be provided.
BRIEF DESCRIPTION OF THE DRAWING
[0022] FIG. 1 is a schematic view of a 180-degree bent specimen
subjected to stress by bolting.
REFERENCE SIGNS LIST
[0023] 1 specimen [0024] 2 bolt
DETAILED DESCRIPTION
[0025] An ultra-high-strength cold-rolled steel sheet has a
specific chemical composition and a microstructure as described
below. The chemical composition of the cold-rolled steel sheet is
first described.
(C: 0.15% to 0.25% by miss)
[0026] C is an element which stabilizes austenite and is necessary
to ensure the strength of the steel sheet. When the content of C is
less than 0.15% by mass, it is difficult for a microstructure
having a tempered martensite phase and a ferrite phase to stably
obtain a tensile strength of 1320 MPa or more. However, when the
content of C is more than 0.25% by mass, welded portions and
heat-affected zones affected by welding are significantly hardened
and therefore weldability is reduced. Therefore, the content of C
is preferably 0.15% to 0.25% by mass and more preferably 0.18% to
0.22% by mass.
(Si: 1.0% to 3,0% by mass)
[0027] Si is a substitutional solid solution hardening element
effective in hardening the steel sheet. The content of Si needs to
be 1.0% by mass or more to develop this effect. When the content of
Si is more than 3.0% by mass, scales are significantly formed
during hot rolling and the failure rate of final products is
increased, which is not economically preferred. Therefore, the
content of Si is 1.0% to 3.0% by mass.
(Mn: 1.5% to 2.5% by mass)
[0028] Mn is an element which stabilizes austertite and is
effective in hardening steel. When the content of Mn is less than
1.5% by mass, it is difficult to stably manufacture the steel sheet
having a target strength because the hardenability of steel is
insufficient, the production of a ferrite phase during cooling from
the annealing temperature and the formation of pearlite and bainite
begin early, and the strength is significantly reduced. However,
when the content thereof is more than 2.5% by mass, segregation is
serious, workability is deteriorated in some cases, and delayed
fracture resistance is reduced. Therefore, the content of Mn is
preferably 1.5% to 2.5% by mass and more preferably 1.5% to 2.0% by
mass.
(P: 0.05% by mass or less)
[0029] P is an element conductive to grain boundary fracture due to
grain boundary segregation and Therefore is preferably low. The
upper limit thereof is 0.05% by mass and is preferably 0.010% by
mass. In view of an increase in weldability. the upper limit
thereof is more preferably 0.008% by mass or less.
(S: 0.02% by mass or less)
[0030] S forms an inclusion such as MnS, causing a reduction in
impact resistance and/or delayed fracture resistance and is
preferably minimized. The upper limit thereof is 0.02% by mass and
preferably 0.002% by mass.
(Al: 0.01% to 0.05% by mass)
[0031] Al is an element effective in deoxidization. The content
thereof needs to be 0,01% by mass or more to achieve an effective
deoxidizing effect. However, when the content thereof is excessive,
more than 0.05% by mass, the steel sheet contains increased amounts
of inclusions and has reduced ductility. Therefore, the content of
Al is 0.01% to 0.05% by mass.
(N: less than 0.005% by mass)
[0032] When the content of N is 0.005% by mass or more, the
formation of nitrides causes a reduction in ductility at high
temperature and low temperature. Therefore, the content of N is
less than 0.005% by mass,
[0033] The steel sheet may further contain one or more of Nb, Ti,
and B as required. The effect of the addition of these three
elements and the preferred content thereof are described below.
(Nb and Ti: 0.1% by mass or less)
[0034] Nb and Ti are elements which have a grain-re:Fining effect
and are effective in increasing the strength of the steel sheet.
Hence, the content of is preferably 0.015% by mass or more.
However, when the content of each of Nb and Ti is more than 0.1% by
mass, the effect thereof is saturated, which is not economically
preferred. Therefore, the content of each of Nb and Ti is 0.1% by
mass or less.
(B: 5 ppm to 30 ppm by mass)
[0035] B is an element effective in increasing the strength of the
steel sheet. The strength-in-creasing effect of B cannot be
expected when the content of B is less than 5 ppm by mass. However,
when the content of B is more than 30 ppm by mass, hot workability
is reduced, which is not preferable in terms of manufacture,
Therefore, the content of B is 5 ppm to 30 ppm by mass.
[0036] The remainders other than the above components are Fe and
unavoidable impurities.
[0037] The microstructure of the cold-rolled steel sheet is
described below.
[0038] We investigated how to increase ductility affecting press
moldability and obtain a steel sheet exhibiting, excellent delayed
fracture resistance after press molding. We found that the
appropriate control of a microstructure is important in exhibiting
high ductility. In particular, it is important that the
microstructure contains 40% or more of a tempered martensite phase
on a volume fraction basis after continuous annealing, the
remainder being a ferrite phase. The microstructure is obtained by
quenching from the annealing temperature and tempering subsequent
to quenching. According to this method, an ultra-high-strength
cold-rolled steel sheet with high ductility can be obtained without
excessively using a transition metal element such as V or Mo,
causing an increase in cost or an alloying element such as Al,
possibly causing casting defects.
[0039] The less the amount of hydrogen permeating through steel is,
the more excellent the delayed fracture resistance is An extremely
large number of dislocations are introduced into the tempered
martensite phase by the crystallographic transformation from an
austenite phase to a martensite phase during quenching. When the
microstructure contains an appropriate amount of the ferrite phase,
the number of the dislocations, which serve as hydrogen-trapping
sites causing delayed fracture, can be more significantly reduced
as compared with a tempered martensite single-phase microstructure
and therefore the amount of hydrogen permeating through can be
reduced.
[0040] The tensile strength of steel with a microstructure
containing a tempered martensite phase and a ferrite phase
increases with an increase in volume fraction of the tempered
martensite phase. This is because the hardness of the tempered
martensite phase is higher than the hardness of the ferrite phase,
the tempered martensite phase, which is a hard phase, exhibits
resistance to deformation during tensile deformation, and the
larger the volume fraction of the tempered martensite phase is, the
more the tensile strength of the steel is close to the tensile
strength of the tempered martensite single-phase microstructure. In
the range of each steel component specified herein a tensile
strength of 1320 MPa or more is not achieved when the volume
fraction of the tempered martensite phase is less than 40%. Since
ductility decreases with at increase in volume fraction of the
tempered martensite phase, a microstructure containing more than
85% of the tempered martensite phase on a volume fraction basis
cannot ensure the volume fraction of the ferrite phase that is
necessary to achieve a high ductility of 12% or more in terms of
total elongation and necessary to increase the delayed fracture
resistance. When the volume fraction of the ferrite phase is less
than 15%, a high ductility of 12% or more in terms of total
elongation is not achieved or an increase in delayed fracture
resistance not sufficient. However, when the volume fraction
thereof is more than 60%, the volume fraction of the tempered
martensite phase that is necessary to achieve a predetermined
strength cannot be ensured.
[0041] From the above reasons, in the microstructure of the
cold-rolled steel sheet, the volume fraction of the tempered
martensite phase and that of the ferrite phase are 40% to 85% and
15% to 60%, respectively, and more preferably 60% to 85% and 15% to
40%, respectively. The microstructure of the cold-rolled steel
sheet may be a two-phase microstructure containing a tempered
martensite phase and ferrite phase each having a desired volume
fraction and may contain a constituent phase such as a retained
austenite phase, a bainite phase, or a pearlite phase, other than
these two phases. However, large amounts of the bainite and
Pearlite phases are present, the bainite phase and the pearlite
phase cause a reduction in ductility and a reduction in strength,
respectively. Therefore, it is not preferable that the
microstructure contains large amounts of the bainite and pearlite
phases. The retained austenite phase is principally present at a
grain boundary in the form of a film, serves as a hydrogen-trapping
site, and therefore may possibly act as an origin of fracture due
to hydrogen embrittlement. Hence, the content thereof is preferably
minimized. Therefore, the volume fraction of the constituent phase
(the retained austenite phase, the bainite phase, or the pearlite
phase) other than the tempered martensite phase and the ferrite
phase is preferably 1% or less in total.
[0042] The tensile strength and ductility (total elongation as
determined by a tensile test using a ES No. 5 tensile specimen) are
1320 MPa or more and 12% or more, respectively. The total
elongation corresponds to the minimum elongation capable of
pressing automobile structural parts such as impact beams. Such a
strength level and elongation level can be readily achieved in our
steel sheets. The delayed fracture resistance is such a performance
that fracture does not occur for 100 hours in a 25.degree. C.
hydrochloric acid environment with a pH of 3. Such a performance
can be readily achieved in our steel sheets.
[0043] Applications of the cold-rolled steel sheet are not
particularly limited. Since the cold-rolled steel sheet has the
above properties, the cold-rolled steel sheet is particularly
suitable for ultra-high-strength automobile safety parts such as
automobile door impact beams and center pillars. Steel sheets
include steel strips. The cold-rolled steel sheet may be subjected
to surface treatment such as plating (electroplating or the like)
or chemical conversion to be used as a surface-treated steel
sheet.
[0044] A method for manufacturing the ultra-high-strength
cold-rolled steel sheet will now be described.
[0045] Steel with the above composition is produced and is then
continuously cast into a cast slab (slab). After being heated to
1200.degree. C. or higher, the slab is hot-rolled at, a finish
rolling end temperature of 800.degree. C. or higher. Reasons for
limiting hot rolling are described below. (Slab-beating temperature
of 1200.degree. C. or higher)
[0046] An increase in rolling load increases the risk of causing
troubles during hot rolling when the heating temperature of the
slab is lower than 1200.degree. C. Thus, the heating temperature of
the slab is 1200.degree. C. or higher. An increase in oxidation
causes an increase in scale loss when the heating temperature
thereof is excessively high. Thus, the heating temperature of the
slab is preferably 1300.degree. C. or lower.
(Finish Rolling End Temperature of 800.degree. C. or Higher)
[0047] A uniform hot-rolled microstructure can be obtained when the
finish rolling end temperature is 800.degree. C. or higher. When
the finish rolling end temperature is lower than 800.degree. C.,
the microstructure of the steel sheet is nonuniform, the ductility
thereof is, and the risk of causing various failures during molding
is increased. Thus, the finish rolling end temperature is
800.degree. C. or higher. The upper limit of the finish rolling end
temperature is not particularly limited and is preferably
1000.degree. C. or lower because rolling at excessively high
temperature causes scale defects.
[0048] The hot-rolled steel sheet is coiled. The coiling
temperature thereof is not particularly limited. When the coiling
temperature thereof is excessively high, the microstructure of the
steel sheet is nonuniform and the ductility thereof is low, due to
formation of coarse grains. When the coiling temperature thereof is
excessively low, a detbrmed microstructure caused by hot rolling
remains to increase the rolling load in cold rolling subsequent to
hot rolling. Therefore, the coiling temperature thereof is
preferably 600.degree. C. to 700.degree. C. In particular, the
coiling temperature thereof is preferably 600.degree. C. to
650.degree. C.
[0049] The hot-rolled steel sheet is pickled, cold-rolled,
continuously annealed, and then tempered. Pickling and cold rolling
conditions are not particularly limited. steel sheet is
continuously annealed such that the steel sheet is held at a
temperature ranging from the Ac.sub.1 transformation temperature to
Ac.sub.3 transformation temperature thereof for 30 s to 1200 s,
cooled to a temperature of 600.degree. C. to 800.degree. C. at an
average cooling rate of 100.degree. C./s or less, and then cooled
to 100.degree. C. or lower at an average cooling rate of
100.degree. C./s to 1000.degree. C./s. The steel sheet is
subsequently tempered such that the steel sheet is reheated and
held at a temperature of 100.degree. C. to 300.degree. C. for 120 s
to 1800 s. Reasons for limiting continuous annealing and tempering
conditions are described below.
(Annealing Temperature: Holding at Temperature Ranging from
Ac.sub.1 Transformation Temperature to Ac.sub.3 Transformation
Temperature for 30 s to 1200 s)
[0050] When the annealing temperature is lower than the Ac.sub.i
transformation temperature, an austenite phase (transformed into a
martensite phase after quenching) necessary to ensure a
predetermined strength is not produced during annealing and
therefore such a predetermined strength cannot be achieved even if
quenching is performed subsequently to annealing. Even if the
annealing temperature is higher than the Ac.sub.3 transformation
temperature, 40% or more of the martensite phase can be obtained on
a volume fraction basis by controlling a ferrite phase precipitated
during cooling from the annealing temperature. In the ease of
performing annealing at a temperature higher than the Ac.sub.3
transformation temperature, a desired microstructure is unlikely to
be obtained. Therefore, the annealing temperature ranges from the
Ac.sub.1 transformation temperature to the Ac.sub.3 transformation
temperature. In view of stably ensuring the equilibrium volume
fraction of the austenite phase to be 40% or more within this
temperature range, 760.degree. C. or higher is preferred and
780.degree. C. or higher is more preferred. When the holding time
(annealing time) at the annealing temperature is excessively short,
a microstructure is not sufficiently annealed, a nonuniform
microstructure in which a deformed microstructure caused by hot
rolling is present is caused, and the ductility is reduced.
However, when the holding time is excessively long, an increase in
manufacturing time is caused, which is not preferable in terms of
manufacturing costs. Therefore, the holding time is 30 seconds to
1200 seconds. In particular, the holding time is preferably 250
seconds to 600 seconds.
(Cooling (Annealing) to Temperature of 600.degree. C. to
800.degree. C. at Average Cooling Rate of 100.degree. C./s or
Less)
[0051] The steel sheet is cooled (the term "cool" is hereinafter
referred to as "anneal" in some cases) to a temperature (annealing
end temperature) of 600.degree. C. to 800.degree. C. from the
annealing temperature at an average cooling rate of 100.degree.
C./s or less. The ferrite phase is precipitated during annealing
from the annealing temperature and the strength-ductility balance
ran be thereby controlled. When the annealing end temperature is
lower than 600.degree. C., a large amount of pearlite is formed in
the microstructure to cause a significant reduction in strength and
therefore a tensile strength of 1320 MPa cannot be achieved. A
sufficient amount of the ferrite phase cannot be precipitated
during annealing from the annealing temperature and therefore
sufficient ductility cannot be achieved when the annealing and
temperature is higher than 800.degree. C. Therefore, the annealing
end temperature is 600.degree. C. to 800.degree. C. The annealing
end temperature is preferably 700.degree. C. to 750.degree. C. to
suppress a change in material quality due to an operational change
in annealing end temperature.
[0052] When the average annealing rate during annealing is more
than 100.degree. C./s a sufficient amount of the ferrite phase is
not precipitated and therefore predetermined ductility cannot be
achieved. The ductility of the microstructure, which contains the
tempered martensite phase and the ferrite phase results from high
work hardenability developed by the coexistence of the tempered
martensite phase, which is hard, and the ferrite phase, which is
soft. The concentration of carbon in the austenite phase during
annealing is insufficient and therefore a hard martensite phase
cannot be obtained during quenching when the average annealing rate
is more than 100.degree. C./s. As a result, the work hardenability
of a final microstructure is reduced and therefore sufficient
ductility is not achieved. Therefore, the average annealing rate
during annealing is 100.degree. C./s or less. The average annealing
rate is preferably 5.degree. C./s or less to sufficiently
concentrate carbon in the austenitic phase.
(Cooling (Quenching) to 100.degree. C. or Lower at Average Cooling
Rate of 100.degree. C./s to 1000.degree. C./s)
[0053] Subsequently to annealing, the steel sheet is cooled (the
term "cool" is hereinafter referred to as "quench" in some cases)
to a temperature (cooling end temperature) of 100.degree. C. or
lower at an average cooling rate of 100.degree. C./s to
1000.degree. C./s. Quenching subsequent to annealing is performed
for the purpose of transforming the austenite phase into the
martensite phase. When the average cooling rate is less than
100.degree. C./s, the austenite phase is transformed into the
ferrite phase, a bainite phase, or a pearlite phase during cooling
and therefore a predetermined strength cannot be achieved. However,
when the average cooling rate is more than 1000.degree. C./s,
shrinkage cracks may possibly be induced in the steel sheet by
cooling. Therefore, the average cooling rate during quenching is
100.degree. C./s to 1000.degree. C./s. The steel sheet is
preferably cooled by water quenching.
[0054] The cooling end temperature is preferably 100.degree. C. or
lower, When the cooling end temperature is higher than 100.degree.
C., the volume fraction of the martensite phase is reduced because
of the insufficient transformation of austenite phase into
martensite phase during quenching and a reduction in material
strength is caused by the self-tempering of the martensite phase
produced by quenching, which is not preferable in terms of
manufacture.
(Tempering: Holding at Temperature of 100.degree. C. to 300.degree.
C. for 120 Seconds to 1800 Seconds)
[0055] Subsequent to quenching, the steel sheet is tempered for the
purpose of tempering the martensite phase such that the steel sheet
is reheated and then held at a temperature of 100.degree. C. to
100.degree. C. for 120 seconds to 1800 seconds. The tempering
thereof softens the martensite phase to increase the workability.
The softening of martensite is insufficient and therefore the
effect of increasing the workability cannot he expected when
performing tempering at lower than 100.degree. C. Performing
tempering at higher than 300.degree. C. increases manufacturing
costs for reheating causes a significant reduction in strength, and
is incapable of achieving a useful effect.
[0056] When the holding time is less than 120 s, martensite phase
is not sufficiently softened at a holding temperature and therefore
the effect of increasing the workability cannot be expected. When
the holding, time is more than 1800 s, the strength is
significantly reduced because of the excessive softening of
martensite phase and manufacturing costs are increased because of
an increase in reheating time, which is not preferable.
[0057] The ultra-high-strength cold-rolled steel sheet can he
manufactured through the above manufacturing steps. Since the
ultra-high-strength cold-rolled steel sheet has excellent
shapeahility (flatness) after annealing, a step of correcting the
shape of the steel sheet by roiling, leveling, or the like is not
necessarily needed. In view of adjusting the quality and/or surface
roughness thereof, the annealed steel sheet may he rolled with an
elongation of several percent.
EXAMPLES
[0058] Test Steels A to M with compositions shown in Table 1 were
produced in a \laMUM and were then formed into slabs, which were
hot-rolled under conditions shown in Table 2, whereby hot-roiled
steel sheets with a thickness of 3.4 mm were prepared. The
hot-rolled steel sheets were surface-desealed by pickling and were
then cold-rolled to a thickness of 1.4 mm. The cold-rolled steel
sheets were continuously annealed and tempered under conditions
shown in Table 2. The Ac.sub.1 transfbrmation temperature and
Ac.sub.3 transformation temperature of each steels is determined
from relational equations (the following two equations) described
in The Japan Institute of Metals, Tekkou Zairyou, Maruzen, 1985, p.
43 and Kinzoku Netsushori Gijutsu Binran Henshuu Enkal, Kinzoku
Netsushori Gijutsu 3rd Edition, The Nikkan Kogyo Shimbun, Ltd.,
1966, p. 137, the equations being involved in the dependence of
transformation temperature on alloying components:
Ac.sub.1(.degree. C.)=723-10.7.times.(% by mass Mn)+29.1.times.(%
by mass Si) (1)
Ac.sub.3(.degree. C.)=910-203.times.(% by mass
C).sup.1/2+29.1.times.(% by mass Si)-30.times.(% by mass
Mn)+700.times.(% by mass P)+400.times.(% by mass Al)+400.times.(%
by mass Ti) (2).
TABLE-US-00001 TABLE 1 A.sub.C1 A.sub.C3 transformation
transformation Steel Composition (% by mass) temperature
temperature symbol C Si Mn P S Al N Ti Nb B (.degree. C.) (.degree.
C.) Remarks A 0.15 1.48 1.8 0.007 0.0011 0.028 0.0031 -- -- -- 747
837 Inventive steel B 0.18 1.48 1.8 0.007 0.0008 0.031 0.0036 -- --
-- 747 830 Inventive steel C 0.25 1.49 1.8 0.010 0.0014 0.027
0.0024 -- -- -- 747 816 Inventive steel D 0.20 1.03 1.8 0.011
0.0008 0.027 0.0027 -- -- -- 734 814 Inventive steel E 0.18 2.97
1.8 0.010 0.0009 0.025 0.0028 -- -- -- 790 873 Inventive steel F
0.20 1.52 1.5 0.011 0.0007 0.033 0.0028 -- -- -- 751 839 Inventive
steel G 0.19 1.54 2.4 0.009 0.0018 0.024 0.0033 -- -- -- 742 810
Inventive steel H 0.18 1.51 1.8 0.010 0.0009 0.026 0.0036 0.04 --
-- 748 847 Inventive steel I 0.18 1.50 1.8 0.010 0.0010 0.038
0.0035 -- 0.04 -- 747 836 Inventive steel J 0.19 1.49 1.8 0.009
0.0010 0.033 0.0029 -- -- 0.002 747 830 Inventive steel K 0.19 1.48
1.8 0.007 0.0012 0.035 0.0037 0.04 0.04 0.002 747 845 Inventive
steel L 0.12 1.46 2.0 0.011 0.0011 0.029 0.0041 -- -- -- 744 841
Comparative steel M 0.15 0.44 1.6 0.009 0.0010 0.022 0.0038 -- --
-- 719 811 Comparative steel
TABLE-US-00002 TABLE 2 Hot rolling step Annealing step Finish
Annealing Annealing Slab-heating rolling Coiling Annealing Holding
average end Steel temperature temperature temperature temperature
time cooling rate temperature No. symbol (.degree. C.) (.degree.
C.) (.degree. C.) (.degree. C.) (s) (.degree. C./s) (.degree. C.) 1
A 1250 900 650 830 600 5 750 2 B 1250 900 650 800 600 14 690 3 B
1250 900 650 800 600 14 710 4 B 1250 900 650 800 600 5 750 5 B 1250
900 650 830 600 19 700 6 C 1250 900 650 800 600 22 650 7 C 1250 900
650 800 600 4 680 8 D 1250 900 650 800 600 5 700 9 E 1250 900 650
800 600 15 750 10 E 1250 900 650 800 600 13 750 11 F 1250 900 650
800 600 5 750 12 G 1250 900 650 780 600 4 620 13 H 1250 900 650 800
600 5 700 14 H 1250 900 650 800 600 12 730 15 I 1250 900 650 800
600 4 680 16 I 1250 900 650 800 600 5 710 17 J 1250 900 650 800 600
3 720 18 K 1250 900 650 800 600 5 720 19 K 1250 900 650 800 600 19
740 20 B 1250 900 650 800 30 12 700 21 B 1250 900 650 800 1200 5
700 22 F 1250 900 650 800 600 4 750 23 F 1250 900 650 800 600 4 750
24 E 1250 900 650 800 10 14 700 25 K 1250 900 650 900 600 24 800 26
L 1250 900 650 830 600 19 650 27 M 1250 900 650 800 600 11 700 28 M
1250 900 650 780 600 10 700 29 M 1250 900 650 780 600 12 750 30 A
1250 900 650 830 600 16 500 31 A 1250 900 650 780 600 20 750 32 B
1250 900 650 830 600 19 700 Annealing step Quenching Tempering step
average Cooling end Tempering Holding cooling rate temperature
temperature time No. (.degree. C./s) (.degree. C.) (.degree. C.)
(s) Remarks 1 904 25 150 1200 Example 2 751 24 150 1200 Example 3
814 22 200 1200 Example 4 973 31 300 1200 Example 5 883 28 300 1200
Example 6 885 25 200 1200 Example 7 833 20 200 1200 Example 8 910
22 200 1200 Example 9 837 21 150 1200 Example 10 767 21 200 1200
Example 11 681 22 150 1200 Example 12 753 24 200 1200 Example 13
625 19 150 1200 Example 14 869 19 200 1200 Example 15 867 23 150
1200 Example 16 855 22 200 1200 Example 17 774 22 200 1200 Example
18 864 23 150 1200 Example 19 995 23 200 1200 Example 20 887 21 200
1200 Example 21 646 20 200 1200 Example 22 984 25 150 150 Example
23 738 24 150 1800 Example 24 846 23 300 1200 Comparative Example
25 967 22 300 1200 Comparative Example 26 941 19 300 1200
Comparative Example 27 910 25 200 1200 Comparative Example 28 809
25 200 1200 Comparative Example 29 811 26 200 1200 Comparative
Example 30 786 20 200 1200 Comparative Example 31 20 19 200 1200
Comparative Example 32 889 20 400 120 Comparative Example
[0059] Specimens were taken from the steel sheets obtained through
the above manufacturing steps, were observed (measured) for
microstructure, and subjected to a tensile test. Furthermore, some
of the steels were subjected to a delayed fracture test. The
results are shown in Table 3.
[0060] The observation (measurement) of microstructure and
performance tests were conducted as described below.
(1) Observation of Microstructure
[0061] Specimens were taken from the obtained cold-rolled steel
sheets. A surface of each specimen that was parallel to the rolling
direction was mirror-polished and was etched with nitai. The
microstructure thereof was observed and photographed with an
optical microscope or a scanning electron microscope, whereby the
type of a constituent phase such as a tempered martensite phase or
a ferrite phase was identify. A photograph of the microstructure
was binarized, whereby the volume fraction of each of the tempered
martensite phase the ferrite phase was determined. Since there was
a possibility that a retained austenite phase was present in the
obtained cold-rolled steel sheets, attempts were made to measure
our examples for retained austenite phase by X-ray (Mo-Ka)
determination. However, the amount of the retained austenite phase
present therein was substantially zero and therefore was not
included in the remainder shown in Table 3.
(2) Tensile test
[0062] JIS No. 5 tensile specimens were taken from the obtained
cold-rolled steel sheets in a direction perpendicular to the
rolling direction and were subjected to a tensile test according to
SIS Z 2241, whereby the specimens were determined for tensile
property (0.2% proof stress (YS)), tensile strength (TS), and total
elongation (EL).
(3) Delayed Fracture Characterization Test
[0063] A specimen with a size of 30 mm.times.100 mm was cut out of
each of the obtained cold-roiled steel sheets such that the
longitudinal direction of the specimen corresponded to the rolling
direction of the cold-rolled steel sheets. An end surface of the
specimen was ground. The specimen was bent to 180 degrees using a
punch having a tip with a radius of curvature of 10 mm. As shown in
FIG. 1, the springback caused in the bent specimen was retained
with a bolt 2 such that the distance between inner portions of the
specimen I was 20 mm. After the specimen i was stressed, the
specimen 1 was immersed in hydrochloric acid with a pH of 3 at
25.degree. C. and was measured for up to 100 hours until the
specimen I was broken. A specimen that was not broken within 100
hours was judged to be acceptable.
TABLE-US-00003 TABLE 3 Volume fraction of Volume Results of
tempered fraction Other delayed fracture Steel martensite of
ferrite constituent YS TS EL characterization No. symbol (%) (%)
phase (MPa) (MPa) (%) test Remarks 1 A 85 15 -- 972 1349 14
Acceptable Example 2 B 64 36 -- 878 1338 14 Acceptable Example 3 B
74 26 -- 1024 1363 15 Acceptable Example 4 B 81 19 -- 1135 1326 14
Acceptable Example 5 B 85 15 -- 1166 1359 12 Acceptable Example 6 C
60 40 -- 819 1378 16 Acceptable Example 7 C 63 37 -- 798 1336 17
Acceptable Example 8 D 81 19 -- 1037 1347 13 Acceptable Example 9 E
45 55 -- 789 1361 19 Acceptable Example 10 E 45 55 -- 863 1332 19
Acceptable Example 11 F 68 32 -- 911 1341 14 Acceptable Example 12
G 60 40 -- 953 1401 13 Acceptable Example 13 H 61 39 -- 819 1327 15
Acceptable Example 14 H 72 28 -- 1092 1394 14 Acceptable Example 15
I 65 35 -- 789 1325 13 Acceptable Example 16 I 74 26 -- 1005 1368
15 Acceptable Example 17 J 80 20 -- 1080 1384 13 Acceptable Example
18 K 60 40 -- 796 1323 16 Acceptable Example 19 K 69 31 -- 1008
1328 15 Acceptable Example 20 B 71 29 -- 997 1349 14 Acceptable
Example 21 B 75 25 -- 976 1354 14 Acceptable Example 22 F 69 31 --
941 1364 14 Acceptable Example 23 F 70 30 -- 1012 1322 13
Acceptable Example 24 E 32 58 Pearlite 951 1068 9 -- Comparative
Example 25 K 100 0 -- 1348 1498 7 Unacceptable Comparative Example
26 L 52 48 -- 889 1075 19 -- Comparative Example 27 M 42 58 -- 632
993 13 -- Comparative Example 28 M 48 52 -- 668 991 14 --
Comparative Example 29 M 100 0 -- 1136 1352 6 Unacceptable
Comparative Example 30 A 24 66 Pearlite 437 658 30 -- Comparative
Example 31 A 0 72 Pearlite 462 578 32 -- Comparative Example 32 B
72 28 -- 983 1166 14 -- Comparative Example
[0064] Tables 1 to 3 confirm that our Examples meet requirements
specified herein and have a tensile strength of 1320 MPa or more, a
total elongation of 12% or more, a high strength-ductility balance,
and excellent delayed fracture resistance because the Examples were
not broken for 100 hours in the delayed fracture characterization
test
[0065] No. 24, of which the annealing time is 10 seconds and
therefore is outside our range, has no predetermined strength or
ductility because a Pearlite phase produced after hot rolling
remains after annealing and the influence of strain due to cold
rolling is not sufficiently removed. Nos. 25 and 29, each of which
the annealing temperature is not lower than the Ac.sub.3
temperature, cannot precipitate any ferrite phase during annealing,
have a martensite single-phase microstructure, and exhibit
predetermined strength and no predetermined ductility, Nos. 26 and
27, of which steel components are outside our range, have no
predetermined strength although continuous annealing and tempering
were performed as specified herein. No. 30, of which the annealing
end temperature is 500.degree. C., contains a large amount of a
ferrite phase precipitated therein and a pearlite phase and
therefore has no predetermined strength. No. 31. of which the
average cooling rate in a quenching step is 20.degree. C./s and
therefore is outside our range, cannot obtain a predetermined
amount of a martensite phase and has no predetermined strength. No.
32, of which the tempering temperature is 400.degree. C., has no
predetermined strength because a martensite phase was excessively
softened by tempering.
[0066] Example Nos. 1 to 23, which meet the requirements specified
herein, were not broken for 100 hours in the delayed fracture
characterization test. This confirms that our cold-rolled steel
sheets have sufficient delayed fracture resistance, However,
Comparative Example Nos. 25 and 29, each of which the
microstructure is a tempered martensite single-phase which is
outside our range, were broken within 100 hours and therefore
failed in the delayed fracture characterization test.
INDUSTRIAL APPLICABILITY
[0067] We provide a thin steel sheet for quenching or tempering,
the thin steel sheet being suitable for use principally in
ultra-high-strength automobile structural parts such as door impact
beams and center pillars for automobiles. In advance of
manufacturing automobile parts from the steel sheet, the
composition, rolling conditions, and annealing conditions are
appropriately controlled. This allows the stud sheet to have a
microstructure containing 40% to 85% of a tempered martensite phase
and 15% to 60% of a ferrite phase on a volume fraction basis, a
tensile strength of 1320 NIPa or more, a total elongation of 12% or
more, an excellent strength ductility balance, and excellent
delayed fracture resistance. The use of an ultra-high-strength
cold-rolled steel sheet enables the pressing of automobile safety
parts such as impact beams, The automobile safety parts exhibit
excellent delayed fracture resistance,
* * * * *