TECHNICAL FIELD
[0001] This invention relates to a high-strength dual-phase steel sheet having an excellent
deep drawability, and particularly to a high-strength dual-phase cold rolled steel
sheet having an excellent deep drawability and a high strength dual phase galvanized
steel sheet having an excellent deep drawability which have a tensile strength of
440 MPa or more and are suitable for use in steel sheets for vehicles as well as a
method of producing the same.
BACKGROUND ART
[0002] Recently, it is required to improve a fuel consumption in a vehicle from a viewpoint
of the maintenance of the global environment, and also it is required to improve a
safety of a vehicle body from a viewpoint of the protection of crews during the collision
of the vehicle. To this end, investigations for achieving both the lightening and
strengthening of the vehicle body are positively proceeding.
[0003] In order to simultaneously satisfy the lightening and strengthening of the vehicle
body, it is said that the high-strengthening of raw materials constituting the parts
is effective, and recently, high-strength steel sheets are positively used as a part
of the vehicle.
[0004] Most of the parts for the vehicle body are formed by press working of the steel sheet
as a raw material. To this end, the high-strength steel sheet used is required to
have an excellent press formability. In order to improve the press formability, it
is necessary to have a high Lankford value (r-value), a high ductility (EI) and a
low yield stress (YS) as mechanical properties of the steel sheet.
[0005] However, in general, as the steel sheet becomes highly strengthened, the r-value
and the ductility lower and the press formability is degraded, while the yield stress
rises to degrade the shapability and hence the problem of springback is apt to occur.
[0006] And also, a high corrosion resistance is required according to a position of the
vehicle part to be applied, so that various surface-treated steel sheets having an
excellent corrosion resistance are used as a steel sheet for the vehicle parts up
to now. Among these surface-treated steel sheets, a galvanized steel sheet is manufactured
in a continuous galvanizing equipment conducting recrystallization annealing and galvanizing
at the same line, so that the provision of an excellent corrosion resistance and a
cheap production are possible. And also, an alloyed galvanized steel sheet obtained
by subjecting to a heat treatment after the galvanization is excellent in the weldability
and press formability in addition to the excellent corrosion resistance. Therefore,
they are widely used.
[0007] In order to further advance the lightening and strengthening of the vehicle body,
in addition to the development of the high-strength cold rolled steel sheet having
the excellent press formability, it is desired to develop a high-strength galvanized
steel sheet having an excellent corrosion resistance through the continuous galvanizing
line.
[0008] As a typical example of the high-strength steel sheet having a good press formability
is mentioned a dual-phase steel sheet having a dual-phase microstructure of a soft
ferrite phase and a hard martensite phase. Especially, the dual-phase steel sheet
produced by cooling with a gas jet after the continuous annealing is low in the yield
stress and possesses a high ductility and an excellent baking hardenability. The above
dual-phase steel sheet is generally good in the workability, but has a drawback that
the workability under severer condition is poor and particularly, the r-value is low
and the deep drawability is bad.
[0009] And also, when the galvanization is applied for providing the excellent corrosion
resistance, the continuous galvanizing line is general to set up the annealing equipment
and the plating equipment continuously. To this end, in case of subjecting to the
galvanization, the cooling after the annealing is constrained by a plating temperature
and can not drop down to a temperature lower than the plating temperature at once
and hence the cooling is interrupted. At a result, an average cooling rate necessarily
becomes smaller. Therefore, when the galvanized steel sheet is produced in the continuous
galvanizing line, it is difficult to generate martensite phase produced under a cooling
condition of a large cooling rate into the steel sheet after the galvanization. To
this end, it is generally difficult to produce the high-strength galvanized steel
sheet having a dual-phase microstructure of a ferrite phase and a martensite phase
through the continuous galvanizing line.
[0010] Under such unfavorable conditions, it is attempted to increase the r-value of the
dual-phase steel sheet to improve the deep drawability. For example,
JP-B-55-10650 discloses a technique that a box annealing is carried out at a temperature ranging
from a recrystallization temperature to A
c3 transformation point after the cold rolling and thereafter the continuous annealing
inclusive of quenching and tempering is carried out after the heating to 700-800°C
in order to obtain the mixed microstructure. In this method, however, the quenching
and tempering are carried out during the continuous annealing, so that the yield stress
is high and hence a low yield ratio can not be obtained. The steel sheet having such
a high yield stress is not suitable for the press formability and has a drawback that
the shapability in the pressed parts is bad.
[0011] And also, a method for lowering the high yield stress is disclosed in
JP-A-55-100934. In this method, the box annealing is first carried out in order to obtain a high
r-value, wherein the temperature in the box annealing is made to a two-phase region
of ferrite (α)-austenite (γ) and Mn is enriched from α phase to γ phase during the
soaking. As the Mn enriched phase preferentially becomes γ phase during the continuous
annealing, the dual-phase microstructure is obtained even at a cooling rate as in
the gas jet cooling, and further the yield stress becomes low. In this method, however,
it is required to conduct the box annealing at a relatively high temperature being
the α-γ two-phase region over a long time for enriching Mn, so that there are many
problems in production steps such as a frequent occurrence of adhesion between steel
sheets inside a coil resulted from the thermal expansion in the annealing, an occurrence
of temper color, a lowering of service life in an inner cover for a furnace body and
the like. Therefore, it was difficult to industrially stably produce high-strength
steel sheets possessing a high r-value and a low yield stress up to now.
[0012] In addition,
JP-B-1-35900 discloses a technique wherein the dual-phase cold rolled steel sheet having a very
high r-value and a low yield stress of r-value = 1.61, YS = 224 MPa and TS = 482 MPa
can be produced by cold rolling a steel having a composition of 0.012 mass% C-0.32
mass% Si-0.53 mass% Mn-0.03 mass% P-0.051 mass% Ti, heating to 870°C corresponding
to α-γ two-phase region and thereafter cooling at an average cooling rate of 100°C/s.
However, the high cooling rate of 100°C/s is difficult to attain in the gas jet cooling
usually used in the continuous annealing line or continuous galvanizing line after
the cold rolling, and is required to use an equipment for water-quenching, and also
a problem becomes actual in the surface treatment of the water-quenched steel sheet,
so that there are problems in the production equipment and the materials.
[0013] Furthermore, it is attempted to produce the high-strength dual-phase galvanized steel
sheet. In the past, as the method of producing the high-strength dual-phase galvanized
steel sheet is generally used a method wherein the formation of low-temperature transformation
phase is facilitated by using a steel added with a large amount of an alloying element
such as Cr or Mo for enhancing a hardenability. However, the addition of the large
amount of the alloying element undesirably brings about the rise of the production
cost.
[0014] Moreover, as is disclosed in
JP-B-62-40405 and the like, there is proposed a method of producing the high-strength dual-phase
galvanized steel sheet by defining the cooling rate after the annealing or the plating
in the continuous galvanizing line. However, this method is not actual from the constraint
on the equipment for the continuous galvanizing line and also the steel sheet obtained
by this method is not said to have a sufficient ductility.
EP 0 969 112 discloses a dual-phase high strength steel having excellent dynamic deformation properties
and a method of production thereof.
DISCLOSURE OF THE INVENTION
[0015] It is, therefore, an object of the invention to solve the aforementioned problems
and to provide high-strength dual-phase cold rolled steel sheets having an excellent
deep drawability and high-strength dual-phase galvanized steel sheets having an excellent
deep drawability as well as a method of producing the same.
[0016] Moreover, the term "galvanized steel sheet" used herein means to include a galvanized
steel sheet obtained by subjecting to a galvanization containing aluminum or the like
in addition to zinc and an alloyed galvanized steel sheet obtained by subjecting to
a heat (alloying) treatment for diffusing iron of the matrix steel sheet into the
plated layer after the galvanization.
[0017] In order to achieve the above object, the inventors have made various studies with
respect to an influence of the alloying element upon the microstructure and the recrystallization
texture in the steel sheet.
[0018] As a result, it has been found that by limiting C in a steel slab to a lower content
and rationalizing V content in relation to C content, before the recrystallization
annealing, C in the steel is precipitated as a V carbide to decrease solid-solute
C as far as possible to thereby develop {111} recrystallization texture to obtain
a high r-value and subsequently the V carbide is dissolved by heating to α-γ two-phase
region to enrich C in austenite for easily generating martensite in a subsequent cooling
process, whereby the high-strength dual-phase cold rolled steel sheet and high-strength
dual-phase galvanized steel sheet having a high r-value and an excellent deep drawability
can be produced stably.
[0019] The results of fundamental experiments performed by the inventors will be explained
below.
[0020] In this case, the experiments are performed with respect to a high-strength dual-phase
cold rolled steel sheet of TS: 590 MPa grade and a high-strength dual-phase cold rolled
steel sheet of TS: 780 MPa grade.
[0021] Firstly, the fundamental experiment in the high-strength dual-phase cold rolled steel
sheet of TS: 590 MPa grade is performed under the following conditions. Each of various
sheet bars having a basic composition of C: 0.03 mass%, Si: 0.02 mass%, Mn: 1.7 mass%,
P: 0.01 mass%, S: 0.005 mass%, Al: 0.04 mass% and N: 0.002 mass% and different V contents
by adding V within a range of 0.03-0.55 mass% is heated to 1250°C and soaked, and
then subjected to three-pass rolling at a finisher delivery temperature of 900°C to
obtain a hot rolled steel sheet having a thickness of 4.0 mm.
[0022] In addition, the same procedure as described above is conducted with respect to various
sheet bars having a basic composition of C: 0.03 mass%, Si: 0.02 mass%, Mn: 1.7 mass%,
P: 0.01 mass%, S: 0.005 mass%, Al: 0.04 mass% and N: 0.002 mass% and different values
of (2×Nb [mass%]/93+2×Ti [mass%]/48)/(V [mass%]/51) by adding V, Nb and Ti within
ranges of 0.03-0.04 mass%, 0.01-0.18 mass% and 0.01-0.18 mass%, respectively, so as
to satisfy a relationship of 0.5×C [mass%]/12 ≤ (V [mass%]/51+2×Nb [mass%]/93+2×Ti
[mass%]/48) ≤ 3×C [mass%]/12.
[0023] Moreover, the hot rolled steel sheet after the finish rolling is subjected to a temperature
holding treatment of 650°C × 1 hour as a coiling treatment. Subsequently, the sheet
is subjected to a cold rolling at a rolling reduction of 70% to obtain a cold rolled
steel sheet having a thickness of 1.2 mm. Next, the cold rolled steel sheet is subjected
to a recrystallization annealing at 850°C for 60 seconds and cooled at a cooling rate
of 30°C/s.
[0024] On the other hand, the fundamental experiment in the high-strength dual-phase cold
rolled steel sheet of TS:780 MPa grade is performed under the following conditions.
[0025] Each of various sheet bars having a basic composition of C: 0.04 mass%, Si: 0.70
mass%, Mn: 2.6 mass%, P: 0.04 mass%, S: 0.005 mass%, Al: 0.04 mass% and N: 0.002 mass%
and different values of (2×Nb/93+2×Ti/48)/(V/51) by adding V, Nb and Ti within ranges
of 0.02-0.06 mass%, 0.01-0.12 mass% and 0.01-0.12 mass%, respectively, so as to satisfy
a relationship of 0.5xC [mass%]/12 ≤ (V [mass%]/51+2×Nb [mass%]/93+2×Ti [mass%]/48)
≤ 3×C [mass%]/12 is heated to 1250°C and soaked, and then subjected to three-pass
rolling at a finisher delivery temperature of 900°C to obtain a hot rolled steel sheet
having a thickness of 4.0 mm. Moreover, the sheet after the finish rolling is subjected
to a temperature holding treatment of 650°C × 1 hour as a coiling treatment. Subsequently,
the sheet is subjected to a cold rolling at a rolling reduction of 70% to obtain a
cold rolled steel sheet having a thickness of 1.2 mm. Next, the cold rolled steel
sheet is subjected to a recrystallization annealing at 850°C for 60 seconds and cooled
at a cooling rate of 30°C/s.
[0026] With respect to the thus obtained cold rolled steel sheets is conducted out a tensile
test to investigate tensile properties. The tensile test is carried out by using JIS
No. 5 tensile test piece. The r-value is determined as an average r-value {= (r
L + r
C + 2xr
D)/4} in a rolling direction (r
L), a direction (r
D) inclined at 45 degree with respect to the rolling direction and a direction (r
C) perpendicular (90°) to the rolling direction.
[0027] FIGS. 1a and 1b show an influence of V content in a steel slab upon r-value and yield
ratio of a cold rolled steel sheet (YR = yield stress (YS)/tensile strength (TS)x
100(%)) in cold rolled steel sheets of TS: 590 MPa grade produced by using a steel
slab containing V but not containing Nb and Ti, V. Moreover, an abscissa in FIGS.
1a and 1b is an atomic ratio ((V/51)/(C/12)) of V content to C content, and an ordinate
is r-value in FIG. 1a and yield ratio (YR) in FIG. 1b.
[0028] As seen from FIGS. 1a and 1b, a high r-value and a low yield ratio are obtained by
limiting V content in the steel slab to a range of 0.5-3.0 as the atomic ratio to
C content and it is possible to produce high-strength dual-phase cold rolled steel
sheet having an excellent deep drawability.
[0029] In the steel sheet according to the invention, the inventors found that a high r-value
is obtained because solid-solute C and N are less and {111} recrystallization texture
is strongly developed before the recrystallization annealing. And also, the inventors
found that by annealing at α-γ two-phase region is dissolved V carbide and the solid-solute
C is enriched into austenite phase in large quantity and the austenite can be easily
transformed into martensite in the subsequent cooling process to obtain a dual-phase
microstructure of ferrite and martensite.
[0030] Although Ti and Nb have mainly been used as a carbide forming element in the past,
the inventors paid notice to V having a solubility of carbide higher than those of
Ti and Nb for effectively obtaining the solid-solute C by annealing at a higher temperature
region. That is, it is found that since V carbide easily dissolves as compared with
Ti carbide and Nb carbide in the annealing at a high temperature, a sufficient amount
of solid-solute C for transforming austenite to martensite is obtained by annealing
at the α-γ two-phase region. In addition, it is clear that this phenomenon is most
remarkably generated by V, but the similar result is obtained by adding Nb and Ti
together.
[0031] Although the invention is based on the above knowledge, the following knowledge is
obtained to achieve another invention.
[0032] The inventors compared r-values in the high-strength dual-phase cold rolled steel
sheets of TS: 590 MPa grade and TS: 780 MPa produced by using steel slabs containing
Nb and Ti in addition to V and made clear the followings. FIGS. 2a and 2b show an
influence of V, Nb and Ti contents in the steel slab upon tensile strength (TS) and
Lankford value (r-value) of a cold rolled steel sheet in the cold rolled steel sheets
of TS: 590 MPa grade and TS: 780 MPa grade produced by using the V, Nb and Ti containing
steel slab. Moreover, an abscissa in FIGS. 2a and 2b is an atomic ratio (2×Nb/93+2×Ti/48)/(V/51)
of Nb and Ti contents to V content, and an ordinate is tensile strength (TS) in FIG.
2a and r-value in FIG. 2b.
[0033] According to the above results, in the TS: 780 MPa grade, the high-strengthening
is attempted by large quantities of solid-solution strengthening elements, so that
the r-value is lowered as compared with that of the TS: 590 MPa grade by the increase
of the solid-solute C content or the like. In the TS: 780 MPa grade, however, the
r-value is considerably improved when the value of (2×Nb/93+2×Ti/48)/(V/51) is a range
of not less than 1.5. Such a characteristic in the TS: 780 MPa grade that the r- value
is remarkably improved when the value of (2×Nb/93+2×Ti/48)/(V/51) is a range of not
less than 1.5 is not recognized in the TS: 590 MPa grade.
[0034] Although the detail of causes on the above result is not clear, it is considered
that in the system containing a large amount of an element resulted in the lowering
of the r-value such as solid-solute C or the like as in the TS: 780 MPa grade, Nb
and Ti easily precipitate the solid-solute C and N as a compound as compared with
V and the solid-solute C and N contents after the hot rolling become less to improve
the r-value. Moreover, when the value of (2×Nb/93+2×Ti/48)/(V/51) exceeds 15, TS considerably
lowers, which is unfavorable for obtaining the high-strength dual-phase cold rolled
steel sheet of TS: 780 MPa grade. This is considered due to the fact that as Nb carbide
and Ti are hardly dissolved as compared with V carbide, if the addition quantities
of the Nb and Ti contents are larger than that of the V content, the C content enriched
in austenite phase is largely decreased in the annealing at the α-γ two-phase region
is widely decreased and martensite phase generated after the cooling is softened.
[0035] The invention is accomplished by further examining based on the above knowledge.
The summary of the invention is as follows.
[0036] A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability,
wherein the steel sheet has a composition comprising further not more than 0.3 mass%
in total of one or tow of Nb: more than 0 mass% but not more than 0.3 mass% and Ti:
more than 0 mass% but not more than 0.3 mass% optionally not more than 2.0 mass% in
total of one or two of Cr and Mo; provided that V, Nb, Ti and C satisfy a relationship
represented by the following equation (ii) instead of the equation (i):
and the remainder being Fe and inevitable impurities.
[0037] Moreover, it is preferable that one or two of Nb: 0.001-0.3 mass% and Ti: 0.001-0.3
mass% is not more than 0.3 mass% in total.
[0038] A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability,
wherein the steel sheet comprises C: 0.03-0.08 mass%, Si: 0.1-2.0 mass%, Mn: 1.0-3.0
mass%, P: not more than 0.05 mass% and S: not more than 0.01 mass% and V, Nb and Ti
satisfy a relationship of 1.5 ≤ (2×Nb/93+2×Ti/48)/(V/51) ≤ 15.
[0039] A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to the item (5), wherein the steel sheet has
a composition comprising further not more than 0.3 mass% in total of one or tow of
Nb: more than 0 mass% but not more than 0.3 mass% and Ti: more than 0 mass% but not
more than 0.3 mass% optionally not more than 2.0 mass% in total of one or two of Cr
and Mo; optionally provided that V, Nb, Ti and C satisfy a relationship represented
by the following equation (iv) instead of the equation (iii):
and the remainder being Fe and inevitable impurities.
[0040] Moreover, it-is preferable that one or two of Nb: 0.001-0.3 mass% and Ti: 0.001-0.3
mass% is not more than 0.3 mass% in total.
[0041] A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to the item (6), wherein the steel slab comprises
C: 0.03-0.08 mass%, Si: 0.1-2.0 mass%, Mn: 1.0-3.0 mass%, P: not more than 0.05 mass%
and S: not more than 0.01 mass% and V, Nb and Ti satisfy a relationship of 1.5 ≤ (2×Nb/93+2×Ti/48)/(V/51)
≤ 15.
[0042] A high-strength dual-phase galvanized steel sheet having an excellent deep drawability
comprising a galvanized coating on the steel sheet disclosed in any one of the items
above.
[0043] A method of producing a high-strength dual-phase galvanized steel sheet having an
excellent deep drawability, wherein a galvanization is carried out after the continuous
annealing at a temperature range from a A
C1 transformation point to a A
C3 transformation point in the production method described in any one of the items (5)-(7).
[0044] A method of producing a high-strength dual-phase galvanized steel sheet having an
excellent deep drawability according to the item (10), which further comprising a
continuous annealing step between the cold rolling step and the continuous annealing
step at a temperature range from a A
C1 transformation point to a A
C3 transformation point.
[0045] The cold rolled steel sheet and the galvanized steel sheet according to the invention
are high-strength dual-phase steel sheets having a tensile strength (TS) of not less
than 440 MPa and an excellent deep drawability.
[0046] At first, the reason of limiting the composition in the cold rolled steel sheet and
the galvanized steel sheet according to the invention will be explained below. Moreover,
mass% represents simply as "%".
C: 0.015-0.08%
[0047] C is an element for increasing the strength of the steel sheet and further promoting
the formation of a dual-phase microstructure of ferrite and martensite, and is necessary
to contain not less than 0.015% from a viewpoint of the formation of the dual-phase
microstructure in the invention. Moreover, if it is intended to increase the strength
to TS: not less than 540 MPa and TS: not less than 780 MPa, the C content is not less
than 0.015% and not less than 0.03%, respectively. On the other hand, when the C content
exceeds 0.08%, the development of {111} recrystallization texture is obstructed to
degrade the deep drawability. Therefore, the invention limits the C content to 0.015-0.08%.
When it is particularly required to increase the strength of the steel sheet, it is
preferable to be 0.03-0.08%. Moreover, it is preferable to be not more than 0.05%
from a viewpoint of the deep drawability.
Si: not more than 2.0%
[0048] Although Si is a useful reinforcing element capable of increasing the strength of
the steel sheet without remarkably lowering the ductility of the steel sheet, if the
content exceeds 2.0%, the deterioration of the deep drawability is caused, but also
the surface properties are degraded. Therefore, Si is limited to not more than 2.0%.
Moreover, if it is intended to increase the strength to TS: not less than 780 MPa,
it is preferable to be not less than 0.1% for ensuring the required strength. And
also, it is preferable to be not less than 0.01% for increasing the strength to TS:
not less than 440 MPa which is a main object of the invention.
Mn: 0.5 - 3.0%
[0049] Mn has an action reinforcing the steel and further has an action of lessening a critical
cooling rate for the obtention of the dual-phase microstructure of ferrite and martensite
to promote the formation of the dual-phase microstructure of ferrite and martensite,
so that it is preferable to contain a content in accordance with the cooling rate
after the recrystallization annealing. And also, Mn is an effective element preventing
the hot tearing through S, so that it is preferable to contain an appropriate content
in accordance with S content. However, when the Mn content exceeds 3.0%, the deep
drawability and weldability are degraded. In the invention, therefore, the Mn content
is limited to not more than 3.0%. Moreover, the Mn content is not less than 0.5% for
remarkably developing the above effect, and particularly it is preferable to be not
less than 1.0% for increasing the strength to TS: not less than 780 MPa. And also,
it is preferable to be not less than 0.1% for increasing the strength to TS: not less
than 440 MPa which is a main object of the invention.
P: not more than 0.10%
[0050] P has an action reinforcing the steel and can be contained in a required amount in
accordance with the desired strength. When the P content exceeds 0.10%, the press
formability is degraded. Therefore, the P content is limited to not more than 0.10%.
Moreover, if a more excellent press formability is required, the P content is preferable
to be not more than 0.08%. Furthermore, when large quantities of C, Mn and the like
are contained in order to ensure TS: not less than 780 MPa, the P content is preferable
to be not more than 0.05% in order to prevent the degradation of the weldability.
In addition, if it is intended to increase the strength to TS: not less than 440 MPa,
it is preferable to be not less than 0.001%.
S: not more than 0.02%
[0051] S is existent as an inclusion in the steel sheet and is an element bringing about
the degradation of the ductility and the formability of the steel sheet, particularly
the stretch-flanging property. Therefore, it is preferable to be decreased as far
as possible, and when it is decreased to not more than 0.02%, S does not exert a bad
influence, so that the S content is 0.02% as an upper limit in the invention. Moreover,
when the more excellent stretch-flanging property is required, or when the large quantities
of C, Mn and the like are contained in order to ensure TS: not less than 780 MPa,
if the excellent weldability is required, the S content is preferable to be not more
than 0.01 %, more preferably not more than 0.005%. On the other hand, the S content
is preferable to be not less than 0.0001% considering a cost for the removal of S
in the steelmaking process. -
Al: 0.005-0.20%
[0052] Al is added to the steel as a deoxidizing element and is a useful element for improving
the cleanliness of the steel, but the addition effect is not obtained at less than
0.005%. On the other hand, when it exceeds 0.20%, the more deoxidizing effect is not
obtained and the deep drawability is inversely degraded. Therefore, the Al content
is limited to 0.005-0.20%. Moreover, the invention does not exclude a steelmaking
method through deoxidization other than the Al deoxidization. For example, Ti deoxidization
or Si deoxidization may be conducted. The steel sheets made by these deoxidizing methods
are included within a scope of the invention. In this case, even if Ca, REM and the
like are added to the molten steel, the characteristics of the steel sheet according
to the invention are not obstructed, so that the steel sheet including Ca, REM and
the like is naturally included within the scope of the invention.
N: not more than 0.02%
[0053] N is an element increasing the strength of the steel sheet by the solid-solution
hardening and the strain ageing hardening, but when N content exceeds 0.02%, the nitride
is increased in the steel sheet to remarkably degrade the deep drawability of the
steel sheet. Therefore, the N content is limited to not more than 0.02%. Moreover,
in case of requiring the more improvement of the press formability, the N content
is preferable to be not more than 0.01 %, more preferably not more than 0.004%. In
this case, considering the cost for denitrification in the steelmaking process, the
N content is preferable to be not less than 0.0001%.
V: 0.01-0.5% and 0,5×C/12 ≤ V/51 ≤ 3×C/12
[0054] V is a most important element in the invention. Before the recrystallization, the
solid-solute C is precipitated and fixed as V carbide to develop the {111} recrystallization
texture, whereby a high r-value can be obtained. Moreover, V dissolves the V carbide
in the annealing at α-γ two-phase region to enrich a large quantity of the solid-solute
C in austenite phase, which is easily transformed into martensite at the subsequent
cooling process, whereby the dual-phase steel sheet having a dual-phase microstructure
of ferrite and martensite can be obtained. Such an effect becomes effective when the
V content is not less than 0.01%, more preferably not less than 0.02% and satisfies
0.5×C/12 ≤ V/51 in relation to the C content. On the other hand, when the V content
exceeds 0.5% or when it is V/51 > 3×C/12 in relation to the C content, the dissolution
of the V carbide at the α-γ two-phase region hardly occurs and the dual-phase microstructure
of ferrite and martensite is hardly obtained. Therefore, the V content is limited
to 0.01-0.5% and to 0.5×C/12 ≤ V/51 ≤ 3×C/12. Moreover, V/51 ≤ 2×C/12 is preferable
for obtaining the dual-phase microstructure of ferrite and martensite.
[0055] In addition to the above composition, it is further preferable to contain not more
than 0.3 (mass)% in total of one or two of Nb: more than 0% but not more than 0.3
(mass)% and Ti: more than 0% but not more than 0.3%, and that V, Nb, Ti contents satisfy
0.5×C/12 ≤ (V/51+2×Nb/93+2×Ti/48) ≤ 3×C/12 in relation to the C content in place of
that the V and C content satisfy 0.5×C/12 ≤ V/51 ≤ 3×C/12. Not more than 0.3% in total
of one or tow of Nb: more than 0% but not more than 0.3% and Ti: more than 0% but
not more than 0.3%, and V, Nb, Ti and C satisfy 0.5×C/12 ≤ (V/51+2×Nb/93+2×Ti/48)
≤ 3×C/12
[0056] Nb and Ti are carbide forming elements likewise V and have the same action as V mentioned
above. That is, a high r-value can be obtained by precipitating and fixing the solid-solute
C as Nb and Ti carbides before the recrystallization to develop the {111} recrystallization
texture, and also a dual-phase steel sheet having a dual-phase microstructure of ferrite
and martensite can be obtained by dissolving the Nb and Ti carbides in the annealing
at the α-γ two-phase region to enrich a large quantity of the solid-solute C in austenite
phase and transforming into martensite in the subsequent cooling process. Moreover,
as the above effect of Nb and Ti is considerably small as compared with that of V,
when only Nb and Ti are added to the steel slab without adding V, the deep drawability
aiming at the invention can not be enhanced sufficiently.
[0057] Therefore, it is preferable to add Nb and Ti of more than 0%. More preferably, each
of the Nb and Ti contents is not less than 0.001%. In this case, it is preferable
to satisfy 0.5×C/12 ≤ (V/51+2×Nb/93+2×Ti/48) in relation to the C and V contents for
developing the above effect. On the other hand, when each of Nb and Ti contents or
both in total thereof exceeds 0.3%, or when the Nb and Ti contents satisfy (V/51+2×Nb/93+2×Ti/48)
> 3×C/12 in relation to the C and V contents, the dissolution of the carbide at the
α-γ two-phase region hardly occurs and hence the dual-phase microstructure of ferrite
and martensite is hardly obtained. Therefore, it is preferable that when either Nb
or Ti is merely added, each of the Nb content and the Ti content is within a range
of more than 0% but not more than 0.3%, and when both of Nb and Ti are added together,
the Nb and Ti contents are not more than 0.3% in total and satisfy 0.5×C/12 ≤ (V/51+2×Nb/93+2×Ti/48)
≤ 3×C/12 in relation to the V and C contents.
[0058] On the other hand, if it is intended to increase the strength to TS: not less than
780 MPa, the deep drawability is apt to be easily degraded by the addition of large
quantities of solid-solution strengthening elements such as C, Mn and the like. In
this case, the V, Nb and Ti contents are further desirable to be a range of 1.5 ≤
(2×Nb/93+2×Ti/48)/ (V/51) ≤ 15. The reason why (2×Nb/93+2×Ti/48)/ (V/51) is limited
to not less than 1.5 is considered due to the fact that although the detail of the
cause is not clear, the formation of carbide after the hot rolling is promoted to
decrease the solid-solute C by adding large quantities of Nb and Ti as compared with
V and hence the {111} recrystallization texture is easily developed. Moreover, in
order to ensure the strength of TS: not less than 780 MPa, (2×Nb/93+2×Ti/48)/ (V/51)
is desirable to be a range of not more than 15.
[0059] Furthermore, in addition to the above steel composition, the steel according to the
invention may further comprise one or two of not more than 2.0% in total of one or
two of Cr and Mo;
not more than 2.0% in total of one or two of Cr and Mo
[0060] All of Cr and Mo in the A-group have an action of decreasing the critical cooling
rate for providing the dual-phase microstructure of ferrite and martensite to promote
the formation of the dual-phase microstructure of ferrite and martensite likewise
Mn and can be included, if necessary. The lower limits of the Cr content and Mo content
preferable for obtaining the above effect are Cr: 0.05%, Mn: 0.05%. However, when
one or two of Cr and Mo exceed 2.0% in total, the deep drawability is degraded. To
this end, one or more of Cr and Mo in the are limited to not more than 2.0% in total.
[0061] The reminder other than the above elements is Fe and inevitable impurities. As the
inevitable impurity are mentioned, for example, Sb, Sn, Zn, Co and the like. As acceptable
ranges of their contents are Sb: not more than 0.01%, Sn: not more than 0.1%, Zn:
not more than 0.01% and Co: not more than 0.1%.
[0062] Next, the microstructure of the steel sheet according to the invention will be explained.
[0063] The cold rolled steel sheet according to the invention has a microstructure consisting
of ferrite phase as a primary phase and a secondary phase including not less than
1% of martensite phase at an area ratio with respect to a whole of the microstructure.
[0064] In order to provide the cold rolled steel sheet having a low yield stress (YS), a
high ductility (El) and an excellent deep drawability, it is required to render the
microstructure of the steel sheet according to the invention into a dual-phase microstructure
consisting of a ferrite phase as a primary phase and a secondary phase including a
martensite phase. It is preferable that the ferrite phase as a primary phase is not
less than 80% at an area ratio and hence the secondary phase is not more than 20%.
When the area ratio of the ferrite phase is less than 80%, it is difficult to ensure
the high ductility and the press formability tends to lower. And also, when a good
ductility is required, it is preferable that the ferrite phase is not less than 85%
at the area ratio and hence the secondary phase is not more than 15%. Moreover, in
order to utilize the advantage of the dual-phase microstructure, the ferrite phase
is required to be not more than 99%.
[0065] In the invention, the secondary phase is required to include the martensite phase
at the area ratio of not less than 1% with respect to the whole of the microstructure.
When the martensite is less than 1% at the area ratio, the low yield stress (YS) and
the high ductility (El) can not be satisfied simultaneously. More preferably, the
martensite phase is not less than 3% but not more than 20% at the area ratio. In case
of requiring a good ductility, the martensite phase is preferable to be not more than
15% at the area ratio. Moreover, the secondary phase may be constituted by only the
martensite phase at the area ratio of not less than 1% or by mixed phases of the martensite
phase at the area ratio of not less than 1% and any of a pearlite phase, a bainite
phase and a retained austenite as an additional phase and is not especially limited.
In the latter case, the pearlite phase, the bainite phase and the retained austenite
are preferable to be not more than 50% in total at the area ratio with respect to
the microstructure of the secondary phase in order to more effectively develop the
effect of the martensite phase.
[0066] The cold rolled steel sheet and the galvanized steel sheet having the above microstructure
are steel sheets having a low yield stress, a high ductility and an excellent deep
drawability.
[0067] Next, the method of producing the cold rolled steel sheet and the galvanized steel
sheet according to the invention will be explained.
[0068] The composition of the steel slab used in the production method of the invention
is the same as the compositions of the aforementioned cold rolled steel sheet and
the galvanized steel sheet, so that the explanation on the reason of the limitation
in the steel slab is omitted.
[0069] The cold rolled steel sheet according to the invention is produced by using a steel
slab having a composition of the above range as a starting material and successively
subjecting this starting material to a hot rolling step of subjecting to a hot rolling
to obtain a hot rolled steel sheet, a pickling step of pickling the hot rolled steel
sheet, a cold rolling step of subjecting the hot rolled steel sheet to a cold rolling
to obtain a cold rolled steel sheet, and a recrystallization annealing step of subjecting
the cold rolled steel sheet to a recrystallization annealing to obtain a cold rolled
annealed steel sheet.
[0070] And also, the galvanized steel sheet according to the invention is produced by using
a steel slab having a composition of the above range as a starting material and successively
subjecting this starting material to a hot rolling step of subjecting to a hot rolling
to obtain a hot rolled steel sheet, a pickling step of pickling the hot rolled steel
sheet, a cold rolling step of subjecting the hot rolled steel sheet to a cold rolling
to obtain a cold rolled steel sheet, and a continuous galvanization step of subjecting
the cold rolled steel sheet to a recrystallization annealing and a galvanizing to
obtain a galvanized steel sheet. Furthermore, it is produced by subjecting the cold
rolled steel sheet to a step of annealing and pickling before the continuous galvanization
step, if necessary.
[0071] The steel slab used is preferable to be produced by a continuous casting process
in order to prevent the macro-segregation of the components, but may be produced by
an ingot casting process or a thin slab casting process. Furthermore, in addition
to the conventional process of cooling to a room temperature once after the production
of the steel slab and again heating, energy-saving processes such as a process for
inserting a hot steel slab into a heating furnace without cooling, a process for direct
sending rolling or direct rolling immediately after slight heat-holding and the like
can be applied without problems.
[0072] The above starting material (steel slab) is subjected to the hot rolling step of
forming the hot rolled steel sheet by heating and hot rolling. In the hot rolling
step, there is particularly no problem even in the use of usual rolling conditions
as .long as the hot rolled steel sheet having a desired thickness can be produced.
Moreover, preferable hot rolling conditions are mentioned below for the reference.
Slab heating temperature: not lower than 900°C
[0073] The slab heating temperature is desirable to be made lower as far as possible in
order to improve the deep drawability by coarsening the precipitate to develop the
{111} recrystallization texture. However, when the slab heating temperature is lower
than 900°C, the rolling load increases and the risk of causing troubles in the hot
rolling increases.
[0074] To this end, the slab heating temperature is preferable to be not lower than 900°C.
And also, the upper limit of the slab heating temperature is more preferable to be
1300°C in terms of the lowering of the yield resulted from the increase of scale loss
accompanied with the increase of the oxide weight. Moreover, it goes without saying
that the utilization of a so-called sheet bar heater of heating the sheet bar in the
hot rolling is an effective process from a viewpoint that the slab heating temperature
is lowered and the troubles in the hot rolling are prevented.
Finisher delivery temperature: not lower than 700°C
[0075] The finisher delivery temperature (FDT) is preferable to be not lower than 700°C
in order to obtain a uniform microstructure of the hot rolled parent sheet for providing
an excellent deep drawability after the cold rolling and the recrystallization annealing.
That is, when the finish deformation temperature is lower than 700°C, not only the
microstructure of the hot rolled parent sheet becomes nonuniform, but also the rolling
load in the hot rolling becomes higher and the risk of causing the trouble in the
hot rolling is increased.
Coiling temperature: not more than 800°C
[0076] The coiling temperature is preferable to be not higher than 800°C. That is, when
the coiling temperature exceeds 800°C, the scale increases and the yield tends to
lower due to the scale loss. And also, when the coiling temperature is lower than
200°C, the shape of the steel sheet remarkably is disordered and the risk of causing
problems in the actual use increases, so that the lower limit of the coiling temperature
is more preferable to be 200°C.
[0077] As mentioned above, in the hot rolling step according to the invention, it is preferable
that the steel slab is heated above 900°C, subjected to the hot rolling at the finish
deformation temperature of not lower than 700°C, and coiled at the coiling temperature
of not higher than 800°C.
[0078] Moreover, in the hot rolling step according to the invention, a lubrication rolling
may be conducted in a part of the finish rolling or between passes thereof in order
to reduce the rolling load in the hot rolling. In addition, the application of the
lubrication rolling is effective from a viewpoint of the uniformization of the steel
sheet shape and the homogenization of the material. Also, the coefficient of friction
in the lubrication rolling is preferable to be within a range of 0.10-0.25.
[0079] Further, the hot rolling step is preferable to be a continuous rolling process wherein
the sheet bars located in front and rear are joined to each other and continuously
subjected to the finish rolling.
[0080] The application of the continuous rolling process is desirable from a viewpoint of
the operating stability in the hot rolling.
[0081] Next, the hot rolled steel sheet is subjected to the pickling for the removal of
the scale. The pickling step is sufficient according to the usual manner and it is
preferable to use a treating solution such as hydrochloric acid, sulfuric acid or
the like as a pickling solution.
[0082] Moreover, the cold rolled steel sheet is formed by subjecting the hot rolled steel
sheet to the cold rolling. The cold rolling conditions are not especially limited
as long as the cold rolled steel sheet having desired size and shape can be obtained,
but it is preferable that a rolling reduction in the cold rolling is not less than
40%. When the rolling reduction is less than 40%, the {111} recrystallization texture
is not developed and the excellent deep drawability can not be obtained.
[0083] The cold rolled steel sheet according to the invention is subjected to a recrystallization
annealing in the subsequent recrystallization annealing step to obtain a cold rolled
annealed steel sheet. The recrystallization annealing is carried out in a continuous
annealing line. On the other hand, the galvanized steel sheet according to the invention
is produced by subjecting the cold rolled steel sheet to recrystallization annealing
and galvanizing in the continuous galvanization line after the cold rolling. In this
case, the annealing temperature in the recrystallization annealing is required to
be conducted at a (α+γ) two-phase region within a temperature range from A
C1 transformation point to A
C3 transformation point. This is due to the fact that the annealing is carried out at
(α+γ) two-phase region to dissolve the carbides of V, Ti and Nb to thereby distribute
an amount of solid-solute C sufficient to transform austenite to martensite into the
austenite phase. When the annealing temperature is lower than the A
C1 transformation point, the microstructure is rendered into the ferrite single phase
and the martensite can not be generated, while when it is higher than the A
C3 transformation point, the crystal grains are coarsened and the microstructure is
rendered into the austenite single phase and the {111} recrystallization texture is
not developed and hence the deep drawability is deteriorated remarkably.
[0084] In the cold rolled steel sheet according to the invention, the cooling in the recrystallization
annealing is preferable to be conducted at a cooling rate of not less than 5°C/s in
order to produce the martensite phase to obtain the dual-phase microstructure of ferrite
and martensite.
[0085] On the other hand, in the galvanized steel sheet according to the invention, it is
preferable to quench to a temperature region of 380-530°C after the above recrystallization
annealing. When a stop temperature of the quenching is lower than 380°C, the defective
plating easily occurs, while when it exceeds 530°C, the unevenness easily occurs on
the plated surface. Moreover, the cooling rate is preferable to be not less than 5°C/s
in order to produce the martensite phase to obtain the dual-phase microstructure of
ferrite and martensite. After the above quenching, the galvanization is carried out
by dipping in a galvanizing bath. In this case, Al concentration in the galvanizing
bath is preferable to be within a range of 0.12-0.145 mass%. When the Al concentration
in the galvanizing bath is less than 0.12 mass%, the alloying excessively advances
and the plating adhesion (resistance to powdering) tends to be deteriorated, while
when it exceeds 0.145 mass%, the defective plating easily occurs.
[0086] And also, the plated layer may be subjected to an alloying treatment after the galvanization.
Moreover, the alloying treatment is preferable to be conducted so that Fe content
in the plated layer is 9-12%.
[0087] As the alloying treatment, it is preferable to conduct the alloying of the galvanized
layer by reheating up to a temperature region of 450-550°C. After the alloying treatment,
it is preferable to cool at a cooling rate of not less than 5°C/s to 300°C. The alloying
at a high temperature is difficult to form the martensite phase and there is caused
a fear of degrading the ductility of the steel sheet, while when the alloying temperature
is lower than 450°C, the progress of the alloying is slow and the productivity tends
to lower. Furthermore, when the cooling rate after the alloying treatment is extremely
small, the formation of the martensite becomes difficult. To this end, the cooling
rate at a temperature region from after the alloying treatment to 300°C is preferable
to be not less than 5°C/s.
[0088] Moreover, if it is required to further improve the plating property, it is preferable
that after the cold rolling and before being subjected to the continuous galvanization,
the annealing is separately conducted in the continuous annealing line and subsequently
an enriched layer of components in the steel produced on the surface of the steel
sheet is removed by pickling and thereafter the above treatment is conducted in the
continuous galvanization line. In this case, the pickling may be carried out in the
pickling line or in the pickling bath arranged in the continuous galvanization line.
Also, the atmosphere in the continuous annealing line is preferable to be a reducing
atmosphere with respect to the steel sheet in order to prevent the formation of the
scale, and it is generally sufficient to use a nitrogen gas containing several % of
H
2. The annealing is preferable to be conducted under a condition that a temperature
of the steel sheet reaching in the continuous annealing line is not lower than the
A
C1 transformation point decided by the components in the steel. Because it is required
to promote the enrichment of the alloying element on the surface of the steel sheet
and to enrich the alloying element in the secondary phase by once forming the dual-phase
microstructure in the continuous annealing line. In the steel sheet after the annealing
in the continuous annealing line, there is a tendency that P among the components
in the steel is diffused to segregate on the surface of the steel sheet and Si, Mn,
Cr and the like enrich as an oxide, so that it is preferable to remove the enriched
layer formed on the surface of the steel sheet by the pickling. Then, the same annealing
as in the above is performed in the continuous galvanization line. In order to develop
the characteristics as the dual-phase microstructure, the annealing in the continuous
galvanization line is preferable to be performed at (α+γ) two-phase region within
a temperature range of from the A
C1 transformation point to the A
C3 transformation point. In this case, the reason why the annealing is performed at
not lower than the A
C1 transformation point in both the continuous annealing line and the continuous galvanization
line is due to the fact that the dual-phase microstructure is formed as mentioned
above. Once an enriching place of the element as the secondary phase is formed by
forming the dual-phase microstructure as a final microstructure in the continuous
annealing line, it becomes possible to enrich the alloying element to some degree
at this place. Desirably, it is sufficient to obtain the same dual-phase microstructure
as in a final product after the cooling, so that the alloying element is more preferable
to be enriched in the vicinity of a triple point of grain boundary (intersection of
the grain boundary formed by three crystal grains). Thereafter, when the annealing
is performed at the two-phase region in the continuous galvanization line, the alloying
element is further enriched in the secondary phase or γ-phase and hence the γ-phase
easily transforms into the martensite phase during the cooling process. Moreover,
the term "alloying element" used herein means a substitutional alloying element such
as Mn, Mo or the like, which makes a situation that diffusion hardly occurs and enrichment
easily occurs at the temperature in the annealing step in order to lower the yield
ratio.
[0089] And also, the cold rolled steel sheet after the recrystallization annealing process
and the galvanized steel sheet after the plating process or after the alloying process
may be subjected to a temper rolling at a rolling reduction of not more than 10% for
correcting the shape and adjusting the surface roughness and the like. Furthermore,
the cold rolled steel sheet according to the invention can be applied as not only
a cold rolled steel sheet for the working but also a blank of a surface treated steel
sheet for the working. As the surface treated steel sheet for the working are mentioned
tin-plated steel sheets, porcelain enamels and so on in addition to the aforementioned
galvanized steel sheets (including alloyed sheets). There is no problem even when
they are subjected to a treatment such as resin or fat coating, various paintings,
electroplating or the like. Moreover, the galvanized steel sheet according to the
invention may be subjected to a special treatment after the galvanization in order
to improve the chemical conversion property, weldability, press formability, corrosion
resistance and the like.
BRIEF DESCRIPTION OF THE DRAWINGS
[0090]
FIG. 1a is a graph showing an influence of V and C contents in steel upon a Lankford
value (r-value).
FIG. 1b is a graph showing an influence of V and C contents in steel upon a yield
ratio (YR = yield stress(YS) / tensile stress(TS) × 100(%)).
FIG. 2a is a graph showing an influence of a relationship among Nb, Ti and V contents
upon a tensile strength (TS) in the high-strength dual-phase cold rolled steel sheets
of TS: 590 MPa grade and TS: 780 MPa grade.
FIG. 2b is a graph showing an influence of a relationship among Nb, Ti and V contents
upon a Lankford value (r-value) in the high-strength dual-phase cold rolled steel
sheets of TS: 590 MPa grade and TS: 780 MPa grade.
BEST MODE FOR CARRYING OUT THE INVENTION
[0091] Each of molten steels having compositions shown in Tables 1-4 is made in a converter
and subjected to a continuous casting process to obtain a slab. In this case, each
of the slabs having the compositions shown in Tables 1 and 2 is prepared for the purpose
of experiments with respect to the cold rolled steel sheet, and each of the slabs
having the compositions shown in Tables 3 and 4 is prepared for the purpose of experiments
with respect to the galvanized steel sheet. Especially, the slabs shown in Tables
2 and 4 are prepared for the purpose of obtaining the cold rolled steel sheet and
galvanized steel sheet of TS: not less than 780 MPa, respectively. Then, the steel
slab is heated to 1150°C and subjected to a hot rolling under conditions of a finish
deformation temperature: 900°C and a coiling temperature: 650°C at a hot rolling step
to obtain a hot rolled steel strip having a thickness of 4.0 mm. Subsequently, the
hot rolled steel strip is pickled and subjected to a cold rolling at a rolling reduction
of 70% at a cold rolling step to obtain a cold rolled steel strip or a cold rolled
sheet having a thickness of 1.2 mm. Next, each of the cold rolled steel sheets in
Tables 1 and 2 is subjected to a recrystallization annealing at an annealing temperature
shown in Tables 5 and 6 in a continuous annealing line. The thus obtained cold rolled
sheet is further subjected to a temper rolling at a rolling reduction of 0.8%. With
respect to the galvanized steel sheets, each of the cold rolled sheets in Tables 3
and 4 is subjected to a recrystallization annealing at an annealing temperature shown
in Tables 7 and 8 and further to a galvanizing in a galvanizing bath having an Al
concentration of 0.13% in a continuous galvanization line. Moreover, with respect
to a part of steel sheets (Steel sheet Nos. 52, 68, 69 and 70 in Table 7), the steel
sheet after the cold rolling is subjected to an annealing at 830°C in a continuous
annealing line and then pickled and annealed and galvanized at a galvanizing bath
temperature of 480°C under an Al concentration in the bath of 0.13% in a continuous
galvanization line and further the thus obtained steel strip (galvanized steel sheet)
is subjected to a temper rolling at a rolling reduction of 0.8%. With respect to the
steel sheets 75 and 77 in Table 7, they are subjected to an alloying treatment at
an alloying temperature of 520°C after the galvanization.
[0092] A test piece is cut out from the obtained steel strip and a microstructure thereof
with respect to a section (C section) perpendicular to the rolling direction is imaged
by using an optical microscope or a scanning electron microscope to measure a structure
ratio of ferrite phase as a primary phase and a kind and a structure ratio of a secondary
phase by using an image analysis device. In this case, a specimen for observing the
microstructure is subjected to a mirror-like polishing and an etching with an alcohol
solution containing 2% HNO
3 and then used for the observation. And also, a tensile test piece of JIS No. 5 is
cut out from the steel strip and subjected to a tensile test according to the definition
of JIS Z 2241 to measure a yield stress (YS), a tensile strength (TS), an elongation
(EI), a yield ratio (YR) and a Lankford value (r-value). These results are shown in
Tables 5-8.
Table 5(a)
Cold rolling |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous annealing line (°C) |
Microstructure |
Mechanical properties of cold rolled steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind*1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
EI (%) |
YR (%) |
r-value |
1 |
1-A |
830 |
92 |
M |
8 |
8 |
330 |
600 |
31 |
55 |
1.8 |
Invention example |
2 |
1-B |
830 |
90 |
M |
10 |
10 |
330 |
610 |
30 |
54 |
1.8 |
Invention example |
3 |
1-B |
980 |
0 |
P, B, M |
15 |
100 |
650 |
720 |
22 |
90 |
0.9 |
Comparative example |
4 |
1-B |
680 |
100 |
- |
0 |
0 |
450 |
530 |
29 |
85 |
0.8 |
Comparative Example |
5 |
1-C |
830 |
92 |
M |
|
8 |
340 |
600 |
31 |
57 |
1.8 |
Invention example |
6 |
1-D |
830 |
90 |
M |
10 |
10 |
330 |
610 |
30 |
54 |
1.4 |
Invention example |
7 |
1-E |
830 |
92 |
M |
8 |
8 |
310 |
570 |
33 |
54 |
1.7 |
Invention example |
8 |
1-F |
830 |
100 |
- |
0 |
0 |
510 |
600 |
27 |
85 |
1.8 |
Comparative example |
9 |
1-G |
830 |
93 |
M |
7 |
7 |
330 |
610 |
31 |
54 |
0.8 |
Comparative example |
10 |
1-H |
850 |
92 |
M |
8 |
8 |
350 |
630 |
29 |
56 |
1.9 |
Invention example |
11 |
1-I |
850 |
93 |
M |
7 |
7 |
330 |
620 |
30 |
53 |
1.9 |
Invention example |
12 |
1-J |
850 |
92 |
M |
8 |
8 |
330 |
610 |
33 |
54 |
1.8 |
Invention example |
13 |
1-K* |
830 |
92 |
M |
8 |
8 |
245 |
450 |
38 |
54 |
1.9 |
Invention example |
14 |
1-L |
830 |
93 |
M |
7 |
7 |
330 |
605 example |
30 |
55 |
1.8 |
Invention |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase.
* out of claimed ranges |
Table 5(b)
Cold rolling |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous annealing line (°C) |
Microstructure |
Mechanical properties of cold rolled steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind*1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
EI (%) |
YR (%) |
r-value |
15 |
1-M |
830 |
92 |
M |
8 |
8 |
340 |
620 |
30 |
55 |
1.7 |
Invention example |
16 |
1-N |
830 |
93 |
M |
7 |
7 |
320 |
600 |
31 |
53 |
1.7 |
Invention example |
17 |
1-O |
830 |
92 |
M, B |
6 |
8 |
340 |
625 |
29 |
54 |
1.8 |
Invention example |
18 |
1-P |
830 |
100 |
- |
0 |
0 |
425 |
520 |
34 |
82 |
1.9 |
Comparative Example |
19 |
1-Q |
830 |
65 |
M |
35 |
35 |
395 |
670 |
29 |
59 |
0.8 |
Comparative example |
20 |
1-R |
850 |
69 |
M |
31 |
31 |
370 |
620 |
30 |
60 |
0.8 |
Comparative example |
21 |
1-S |
850 |
100 |
- |
0 |
0 |
495 |
615 |
30 |
80 |
1.7 |
Comparative example |
22 |
1-T |
850 |
92 |
M |
8 |
8 |
355 |
575 |
32 |
62 |
1.7 |
Invention example |
23 |
1-U |
850 |
100 |
- |
0 |
0 |
470 |
580 |
31 |
81 |
1.8 |
Comparative example |
24 |
1-V |
830 |
91 |
M |
9 |
9 |
350 |
570 |
32 |
61 |
1.7 |
Invention example |
25 |
1-W |
850 |
100 |
- |
0 |
0 |
480 |
595 |
31 |
81 |
1.8 |
Comparative example |
26 |
1-X |
830 |
72 |
M |
28 |
28 |
350 |
560 |
31 |
63 |
0.8 |
Comparative example |
27 |
1-Y |
830 |
100 |
- |
0 |
0 |
475 |
590 |
30 |
81 |
1.7 |
Comparative example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
Table 6(a)
Cold rolling |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous annealing line (°C) |
Microstructure |
Mechanical properties of cold rolled steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind *1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
EI (%) |
YR (%) |
r-value |
28 |
2-A |
780 |
90 |
M |
10 |
10 |
560 |
825 |
19 |
68 |
1.1 |
Invention example |
29 |
2-B |
780 |
87 |
M |
13 |
13 |
550 |
810 |
19 |
68 |
1.3 |
Invention example |
30 |
2-B |
950 |
0 |
P,B,M |
19 |
100 |
740 |
860 |
16 |
86 |
0.7 |
Comparative example |
31 |
2-B |
680 |
100 |
- |
0 |
0 |
625 |
770 |
22 |
81 |
0.8 |
Comparative Example |
32 |
2-C |
750 |
88 |
M |
12 |
12 |
540 |
805 |
20 |
67 |
1.3 |
Invention example |
33 |
2-D |
760 |
88 |
M |
12 |
12 |
545 |
810 |
19 |
67 |
1.2 |
Invention example |
34 |
2-E |
770 |
87 |
M |
13 |
13 |
550 |
820 |
20 |
67 |
1.3 |
Invention example |
35 |
2-F |
780 |
100 |
- |
0 |
0 |
660 |
830 |
19 |
80 |
1.4 |
Comparative example |
36 |
2-G |
780 |
69 |
M |
31 |
31 |
540 |
820 |
20 |
66 |
0.7 |
Comparative example |
37 |
2-H |
760 |
81 |
M |
19 |
19 |
620 |
930 |
15 |
67 |
1.3 |
Invention example |
38 |
2-I |
780 |
83 |
M |
17 |
17 |
590 |
860 |
17 |
69 |
1.1 |
Invention example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
Table 6(b)
Cold rolling |
Steel sheet No. |
Steel No. |
Annealing temperature in No. continuous annealing line (°C) |
Microstructure |
Mechanical properties of cold rolled steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind*1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
EI (%) |
YR (%) |
r-value |
39 |
2-J |
780 |
87 |
M |
13 |
13 |
445 |
660 |
27 |
67 |
1.4 |
Invention example |
40 |
2-K |
760 |
68 |
M |
32 |
32 |
570 |
850 |
18 |
67 |
0.8 |
Comparative example |
41 |
2-L |
780 |
100 |
- |
0 |
0 |
690 |
835 |
19 |
83 |
1.3 |
Comparative example |
42 |
2-M |
780 |
85 |
M |
15 |
15 |
525 |
805 |
20 |
65 |
1.1 |
Invention example |
43 |
2-N |
760 |
88 |
M |
12 |
12 |
530 |
800 |
20 |
66 |
1.3 |
Invention example |
44 |
2-O |
780 |
90 |
M |
10 |
10 |
525 |
790 |
21 |
66 |
1.3 |
Invention example |
45 |
2-P |
780 |
100 |
- |
0 |
0 |
650 |
795 |
21 |
82 |
1.3 |
Comparative example |
46 |
2-Q |
760 |
87 |
M |
13 |
13 |
540 |
810 |
19 |
67 |
1.1 |
Invention example |
47 |
2-R |
760 |
88 |
M |
12 |
12 |
545 |
815 |
15 |
67 |
1.3 |
Invention example |
48 |
2-S |
780 |
90 |
M |
10 |
10 |
540 |
810 |
19 |
67 |
1.3 |
Invention example |
49 |
2-T |
780 |
100 |
- |
0 |
0 |
665 |
785 |
20 |
85 |
1.4 |
Comparative example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
Table 7(a)
Galvanizing |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous galvanization line (°C) |
Microstructure |
Mechanical properties of galvanized steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties Tensile properties |
Area ratio (%) |
Kind *1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
El (%) |
YR (%) |
r-value |
50 |
3-A |
830 |
92 |
M |
8 |
8 |
330 |
610 |
31 |
54 |
1.7 |
Invention example |
51 |
3-B |
830 |
90 |
M |
10 |
10 |
330 |
620 |
30 |
53 |
1.7 |
Invention example |
52 |
3-B |
830 |
92 |
M |
8 |
8 |
350 |
630 |
30 |
56 |
1.6 |
Invention example |
53 |
3-B |
980 |
0 |
P, B, M |
12 |
100 |
660 |
720 |
22 |
92 |
0.9 |
Comparative Example |
54 |
3-B |
680 |
100 |
- |
0 |
0 |
460 |
540 |
28 |
85 |
0.8 |
Comparative example |
55 |
3-C |
830 |
90 |
M |
10 |
10 |
340 |
610 |
31 |
56 |
1.7 |
Invention example |
56 |
3-D |
830 |
92 |
M |
8 |
8 |
340 |
620 |
30 |
55 |
1.4 |
Invention example |
57 |
3-E |
830 |
94 |
M |
6 |
6 |
320 |
580 |
32 |
55 |
1.6 |
Invention example |
58 |
3-F |
830 |
100 |
- |
0 |
0 |
510 |
600 |
27 |
85 |
1.7 |
Comparative example |
59 |
3-G |
830 |
92 |
M |
8 |
8 |
330 |
610 |
30 |
54 |
0.8 |
Comparative example |
60 |
3-H |
850 |
93 |
M |
7 |
7 |
340 |
630 |
30 |
54 |
1.8 |
Invention example |
61 |
3-1 |
850 |
92 |
M |
8 |
8 |
340 |
620 |
31 |
55 |
1.8 |
Invention example |
62 |
3-J |
850 |
92 |
M |
8 |
8 |
320 |
610 |
31 |
52 |
1.7 |
Invention example |
63 |
3-K |
830 |
92 |
M, B |
6 |
8 |
330 |
610 |
30 |
54 |
1.6 |
Invention example |
64 |
3-L* |
830 |
92 |
M |
8 |
8 |
248 |
450 |
37 |
55 |
1.7 |
Invention example |
65 |
3-M |
830 |
93 |
M |
7 |
7 |
340 |
620 |
30 |
55 |
1.6 |
Invention example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase.
*out of claimed ranges |
Table 7(b)
Galvanizing |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous galvanization line (°C) |
Microstructure |
Mechanical properties of galvanized steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind *1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
El (%) |
YR (%) |
r-value |
66 |
3-N |
830 |
92 |
M |
8 |
8 |
320 |
600 |
31 |
53 |
1.6 |
Invention example |
67 |
3-0 |
830 |
93 |
M |
7 |
7 |
340 |
625 |
29 |
54 |
1.7 |
Invention example |
68 |
3-H |
830 |
92 |
M |
8 |
8 |
340 |
620 |
30 |
55 |
1.8 |
Invention example |
69 |
3-K |
830 |
93 |
M |
7 |
7 |
320 |
600 |
31 |
53 |
1.6 |
Invention example |
70 |
3-M |
830 |
92 |
M |
8 |
8 |
320 |
610 |
31 |
52 |
1.6 |
Invention example |
71 |
3-P |
830 |
100 |
- |
0 |
0 |
420 |
510 |
34 |
82 |
1.8 |
Comparative example |
72 |
3-Q |
830 |
66 |
M |
34 |
34 |
390 |
670 |
27 |
58 |
0.8 |
Comparative example |
73 |
3-R |
850 |
68 |
M |
32 |
32 |
385 |
615 |
30 |
63 |
0.8 |
Comparative example |
74 |
3-S |
850 |
100 |
- |
0 |
0 |
500 |
605 |
31 |
83 |
1.6 |
Comparative example |
75 |
3-T |
850 |
91 |
M |
9 |
9 |
350 |
580 |
31 |
60 |
1.7 |
Invention example |
76 |
3-U |
850 |
100 |
- |
0 |
0 |
480 |
575 |
32 |
83 |
1.6 |
Comparative example |
77 |
3-V |
830 |
91 |
M |
9 |
9 |
340 |
580 |
31 |
59 |
1.7 |
Invention example |
78 |
3-W |
850 |
100 |
- |
0 - |
0 |
490 |
600 |
30 |
82 |
1.7 |
Comparative example |
79 |
3-X |
830 |
70 |
M |
30 |
30 |
340 |
565 |
32 |
60 |
0.8 |
Comparative example |
80 |
3-Y |
830 |
100 |
- |
0 |
0 |
490 |
600 |
30 |
82 |
1.7 |
Comparative example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
Table 8(a)
Galvanizing |
Steel sheet No. |
Steel No. |
Annealing temperature in continuous galvanization (°C) |
Microstructure |
Mechanical properties of galvanized steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind *1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (MPa) |
TS (MPa) |
El (%) |
YR |
r-value |
81 |
4-A |
780 |
91 |
M |
9 |
9 |
560 |
815 |
19 |
69 |
1.1 |
Invention example |
82 |
4-B |
780 |
89 |
M |
11 |
11 |
555 |
805 |
19 |
69 |
1.4 |
Invention example |
83 |
4-B |
950 |
0 |
P,B,M |
21 |
100 |
735 |
850 |
16 |
86 |
0.8 |
Comparative example |
84 |
4-B |
680 |
100 |
- |
0 |
0 |
620 |
760 |
22 |
82 |
0.8 |
Comparative Example |
85 |
4-C |
4 |
89 |
M |
11 |
11 |
545 |
800 |
20 |
68 |
1.3 |
Invention example |
86 |
4-D |
760 |
88 |
M |
12 |
12 |
550 |
805 |
19 |
68 |
1.4 |
Invention example |
87 |
4-E |
770 |
90 |
M |
10 |
10 |
550 |
810 |
20 |
68 |
1.3 |
Invention example |
88 |
4-F |
780 |
100 |
- |
0 |
0 |
675 |
815 |
19 |
83 |
1.5 |
Comparative example |
89 |
4-G |
780 |
92 |
M |
8 |
8 |
550 |
810 |
20 |
68 |
0.8 |
Comparative example |
90 |
4-H |
760 |
83 |
M |
17 |
17 |
635 |
935 |
15 |
68 |
1.3 |
Invention example |
91 |
4-I |
780 |
85 |
M |
15 |
15 |
590 |
855 |
17 |
69 |
1.1 |
Invention example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
Table 8(b)
Galvanizing |
Steel sheet No. |
Steel No. |
Annealing temperature continuous galvanization line (°C) |
Microstructure |
Mechanical properties of galvanized steel sheet |
Remarks |
Ferrite phase |
Second phase |
Tensile properties |
Area ratio (%) |
Kind *1 |
Area ratio of martensite (%) |
Area ratio of second phase (%) |
YS (Mpa) |
TS (MPa) |
EI (%) |
YR (%) |
r-value |
92 |
4-J |
780 |
85 |
M |
15 |
15 |
440 |
665 |
25 |
68 |
1.4 |
Invention example |
93 |
4-K |
760 |
67 |
M |
33 |
33 |
560 |
860 |
18 |
65 |
0.8 |
Comparative example |
94 |
4-L |
780 |
100 |
- |
0 |
0 |
695 |
840 |
19 |
83 |
1.4 |
Comparative example |
95 |
4-M |
780 |
86 |
M |
14 |
14 |
510 |
810 |
20 |
63 |
1.1 |
Invention example |
96 |
4-N |
760 |
89 |
M |
11 |
11 |
525 |
800 |
20 |
66 |
1.3 |
Invention example |
97 |
4-O |
780 |
89 |
M |
11 |
11 |
525 |
795 |
20 |
66 |
1.3 |
Invention example |
98 |
4-P |
780 |
100 |
- |
0 |
0 |
660 |
805 |
20 |
82 |
1.4 |
Comparative example |
99 |
4-Q |
760 |
87 |
M |
13 |
13 |
525 |
810 |
19 |
65 |
1.1 |
Invention example |
100 |
4-R |
760 |
86 |
M |
14 |
14 |
530 |
810 |
19 |
65 |
1.2 |
Invention example |
101 |
4-S |
780 |
89 |
M |
11 |
11 |
540 |
820 |
18 |
66 |
1.3 |
Invention example |
102 |
4-T |
780 |
100 |
- |
0 |
0 |
660 |
790 |
20 |
84 |
1.3 |
Comparative example |
(Note) *1: F is abbreviation of ferrite phase, M is abbreviation of matensite phase,
P is abbreviation of perlite phase and B is abbreviation of beinite phase. |
[0093] As seen from the results shown in Tables 5 and 6, the cold rolled steel sheets in
all invention examples have a low yield stress (YS), a high elongation (El) and a
low yield ratio (YR) and further indicate a high r-value and are excellent in the
deep drawability, and have a tensile strength (TS) of not less than 440 MPa. On the
contrary, in the comparative examples being outside the range of the invention, the
yield stress (YS) is high, the elongation (El) is low, or the r-value is low. Particularly,
the somewhat lowering of the r-value accompanied with the high-strengthening is observed
in the high-strength steel sheets of TS: not less than 780 MPa shown in Table 6, for
example, the steel sheet No. 28 produced by using the steel No. 2-A containing V and
no Nb and Ti and the steel sheet No. 38 produced by using the steel No. 2-I containing
V, Nb and Ti and satisfying a relationship of 0.5xC/12 ≤ (V/51+2×Nb/93+2×Ti/48) ≤
3×C/12 but satisfying a relationship of (2×Nb/93+2×Ti/48)/(V/51) < 0.5. On the other
hand, the r-value is improved in the steel sheet Nos. 29, 32, 33 and 34 produced by
using the steel Nos. 2-B, 2-C, 2-D and 2-E containing V, Nb and Ti and satisfying
both relationships of 0.5×C/12 ≤ (V/51+2×Nb/93+2×Ti/48) ≤ 3×C/12 and 1.5 ≤ (2×Nb/93+2×Ti/48)/(V/51)
≤ 15.
[0094] And also, the results obtained with respect to the galvanized steel sheets are shown
in Tables 7 and 8. Even in these galvanized steel sheets, the results similar to those
of the above cold rolled steel sheets are obtained.
[0095] In the steel sheet according to the invention, excellent properties are obtained
even by the production process conducting the galvanization.
INDUSTRIAL APPLICABILITY
[0096] The invention develops an industrially remarkable effect that the high-strength cold
rolled steel sheet and galvanized steel sheet having an excellent deep drawability
can be produced stably. When the cold rolled steel sheet and the galvanized steel
sheet according to the invention are applied to vehicle parts, there are effects that
the press forming is easy and they can sufficiently contribute to reduce the weight
of the vehicle body.
1. A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability,
characterized in that the steel sheet has a composition comprising C: 0.015 - 0.08 mass%, Si: not more
than 2.0 mass%, Mn: 0.5 - 3.0 mass%, P: not more than 0.10 mass%, S: not more than
0.02 mass%, Al: 0.005 - 0.20 mass%, N: not more than 0.02 mass%, V: 0.01 - 0.5 mass%,
optionally not more than 0.3 mass% in total of one or two of Nb: more than 0 mass%
but not more than 0.3 mass%, and Ti: more than 0 mass% but not more than 0.3 mass%,
and optionally not more than 2.0 mass% in total of one or two of Cr and Mo, provided
that V, Nb, Ti and C - if present - satisfy a relationship of 0.5xC/12 ≤ (V/51+2xNb/93+2xTi/48)
≤ 3xC/12, and the remainder being Fe and inevitable impurities, and has a microstructure
consisting of a ferrite phase as a primary phase and a secondary phase including martensite
phase at an area ratio of not less than 1% to a whole of the microstructure.
2. A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability
according to claim 1, wherein the steel sheet further comprises not more than 0.3
mass% in total of one or two of Nb: more than 0 mass% but not more than 0.3 mass%
and Ti: more than 0 mass% but not more than 0.3 mass%.
3. A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability
according to claim 2, wherein the steel sheet comprises not more than 0.3 mass% in
total of one or two of Nb: 0.001 - 0.3 mass% and Ti: 0.001 - 0.3 mass%.
4. A high-strength dual-phase cold rolled steel sheet having an excellent deep drawability
according to claim 2, wherein the steel sheet comprises C: 0.03 - 0.08 mass%, Si:
0.1 - 2.0 mass%, Mn: 1.0 - 3.0 mass%, P: not more than 0.05 mass% and S: not more
than 0.01 mass%, provided that V, Nb and Ti satisfy a relationship of 1.5 ≤ (2xNb/93+2xTi/48)/(V/51)
≤ 15.
5. A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability, which comprises hot rolling a steel slab having a composition
comprising C: 0.015 - 0.08 mass%, Si: not more than 2.0 mass%, Mn: 0.5 - 3.0 mass%,
P: not more than 0.10 mass%, S: not more than 0.02 mass%, Al: 0.005 - 0.20 mass%,
N: not more than 0.02 mass%, V: 0.01 - 0.5 mass%, optionally not more than 0.3 mass%
in total of one or two of Nb: more than 0 mass% but not more than 0.3 mass%, and Ti:
more than 0 mass% but not more than 0.3 mass%, and optionally not more than 2.0 mass%
in total of one or two of Cr and Mo, provided that V, Nb, Ti and C - if present -
satisfy a relationship of 0.5xC/12 ≤ (V/51+2xNb/93+2xTi/48) ≤ 3xC/12, and the remainder
being Fe and inevitable impurities, pickling, cold rolling and then subjecting to
a continuous annealing at a temperature range from a AC1 transformation point to a AC3 transformation point.
6. A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to claim 5, wherein the steel slab further comprises
not more than 0.3 mass% in total of one or two of Nb: more than 0 mass% but not more
than 0.3 mass% and Ti: more than 0 mass% but not more than 0.3 mass%.
7. A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to claim 6, wherein the steel slab comprises
not more than 0.3 mass% in total of one or two of Nb: 0.001 - 0.3 mass% and Ti: 0.001
- 0.3 mass%.
8. A method of producing a high-strength dual-phase cold rolled steel sheet having an
excellent deep drawability according to claim 6, wherein the steel slab comprises
C: 0.03 - 0.08 mass%, Si: 0.1 - 2.0 mass%, Mn: 1.0 - 3.0 mass%, P: not more than 0.05
mass% and S: not more than 0.01 mass%, provided that V, Nb and Ti satisfy a relationship
of 1.5 ≤ (2xNb/93+2xTi/48)/(V/51) ≤ 15.
9. A high-strength dual-phase galvanized steel sheet having an excellent deep drawability
comprising a galvanized coating on the steel sheet as claimed in any one of claims
1 to 4.
10. A method of producing a high-strength dual-phase galvanized steel sheet having an
excellent deep drawability, characterized by subjecting to a galvanization after the continuous annealing at a temperature range
from a AC1 transformation point to a AC3 transformation point in the method claimed in any one of claims 5 to 8.
11. A method of producing a high-strength dual-phase galvanized steel sheet having an
excellent deep drawability according to claim 10, characterized by further comprising a continuous annealing step between the cold rolling step and
the continuous annealing step at a temperature range from a AC1 transformation point to a AC3 transformation point.
1. Hochfestes zweiphasiges kaltgewalztes Stahlblech mit hervorragender Tiefziehfähigkeit,
dadurch gekennzeichnet, dass das Stahlblech eine Zusammensetzung aufweist, die 0,015 - 0,08 Masse-% C, nicht mehr
als 2,0 Masse-% Si, 0,5 - 3,0 Masse-% Mn, nicht mehr als 0,10 Masse-% P, nicht mehr
als 0,02 Masse-% S, 0,005 - 0,20 Masse-% Al, nicht mehr als 0,02 Masse-% N, 0,01 -
0,5 Masse-% V, optional nicht mehr als insgesamt 0,3 Masse-% an Nb: mehr als 0 Masse-%,
jedoch nicht mehr als 0,3 Masse-%, und/oder Ti: mehr als 0 Masse-%, jedoch nicht mehr
als 0,3 Masse-%, und optional nicht mehr als insgesamt 2,0 Mass-% an Cr und/oder Mo
umfasst, mit der Maßgabe, dass V, Nb, Ti und C - falls vorhanden - die Beziehung 0,5C/12
≤ (V/51 + 2xNb/93 + 2Ti/48) ≤ 3xC/12 erfüllen, und wobei der Rest Fe und beiläufige
Verunreinigungen sind, und eine Mikrostruktur aufweist, die aus einer Ferritphase
als Primärphase und einer Sekundärphase, die eine Martensitphase mit einem Flächenanteil
von nicht weniger als 1% an der Gesamtmikrostruktur enthält, besteht.
2. Hochfestes zweiphasiges kaltgewalztes Stahlblech mit hervorragender Tiefziehfähigkeit
nach Anspruch 1, wobei das Stahlblech ferner nicht mehr als insgesamt 0,3 Masse-%
an Nb: mehr als 0 Masse-%, jedoch nicht mehr als 0,3 Masse-%, und/oder Ti: mehr als
0 Masse-%, jedoch nicht mehr als 0,3 Masse-%, umfasst.
3. Hochfestes zweiphasiges kaltgewalztes Stahlblech mit hervorragender Tiefziehfähigkeit
nach Anspruch 2, wobei das Stahlblech nicht mehr als 0,3 Masse-% an Nb: 0,001 - 0,3
Masse-%, und/oder Ti: 0,001 - 0,3 Masse-%, umfasst.
4. Hochfestes zweiphasiges kaltgewalztes Stahlblech mit hervorragender Tiefziehfähigkeit
nach Anspruch 2, wobei das Stahlblech 0,03 - 0,08 Masse-% C, 0,1 - 2,0 Masse-% Si,
1,0 - 3,0 Masse-% Mn, nicht mehr als 0,05 Masse-% P und nicht mehr als 0,01 Masse-%
S umfasst, mit der Maßgabe, dass V, Nb und Ti die Beziehung
erfüllen.
5. Verfahren zur Herstellung eines hochfesten zweiphasigen kaltgewalzten Stahlblechs
mit hervorragender Tiefziehfähigkeit, wobei das Verfahren das Warmwalzen eines Stahlwalzblocks
mit einer Zusammensetzung, die 0,015 - 0,08 Masse-% C, nicht mehr als 2,0 Masse-%
Si, 0,5 - 3,0 Masse-% Mn, nicht mehr als 0,10 Masse-% P, nicht mehr als 0,02 Masse-%
S, 0,005 - 0,20 Masse-% Al, nicht mehr als 0,02 Masse-% N, 0,01 - 0,5 Masse-% V, optional
nicht mehr als insgesamt 0,3 Masse-% an Nb: mehr als 0 Masse-%, jedoch nicht mehr
als 0,3 Masse-%, und/oder Ti: mehr als 0 Masse-%, jedoch nicht mehr als 0,3 Masse-%,
und optional nicht mehr als insgesamt 2,0 Mass-% an Cr und/oder Mo umfasst, mit der
Maßgabe, dass V, Nb, Ti und C - falls vorhanden - die Beziehung 0,5xC/12 ≤ (V/51 +
2xNb/93 + 2xTi/48) ≤ 3xC/12 erfüllen, und wobei der Rest Fe und beiläufige Verunreinigungen
sind, Beizen, Kaltwalzen und dann Durchführen eines kontinuierlichen Glühens in einem
Temperaturbereich vom AC1-Transformationspunkt zum AC3-Transformationspunkt umfasst.
6. Verfahren zur Herstellung eines hochfesten zweiphasigen kaltgewalzten Stahlblechs
mit hervorragender Tiefziehfähigkeit nach Anspruch 5, wobei der Stahlwalzblock ferner
nicht mehr als insgesamt 0,3 Masse-% an Nb: mehr als 0 Masse-%, jedoch nicht mehr
als 0,3 Masse-%, und/oder Ti: mehr als 0 Masse-%, jedoch nicht mehr als 0,3 Masse-%
umfasst.
7. Verfahren zur Herstellung eines hochfesten zweiphasigen kaltgewalzten Stahlblechs
mit hervorragender Tiefziehfähigkeit nach Anspruch 6, wobei der Stahlwalzblock nicht
mehr als insgesamt 0,3 Masse-% an Nb: 0,001 - 0,3 Masse-%, und/oder Ti: 0,001 - 0,3
Masse-%, umfasst.
8. Verfahren zur Herstellung eines hochfesten zweiphasigen kaltgewalzten Stahlblechs
mit hervorragender Tiefziehfähigkeit nach Anspruch 6, wobei der Stahlwalzblock 0,03
- 0,08 Masse-% C, 0,1 - 2,0 Masse-% Si, 1,0 - 3,0 Masse-% Mn, nicht mehr als 0,05
Masse-% P und nicht mehr als 0,01 Masse-% S umfasst, mit der Maßgabe, dass V, Nb und
Ti die Beziehung
erfüllen.
9. Hochfestes zweiphasiges verzinktes Stahlblech mit hervorragender Tiefziehfähigkeit,
das eine Verzinkungsschicht auf dem Stahlblech gemäß einem der Ansprüche 1 bis 4 umfasst.
10. Verfahren zur Herstellung eines hochfesten zweiphasigen verzinkten Stahlblechs mit
hervorragender Tiefziehfähigkeit, dadurch gekennzeichnet, dass in dem Verfahren gemäß einem der Ansprüche 5 bis 8 ein Verzinken nach dem kontinuierlichen
Glühen in einem Temperaturbereich vom AC1-Transformationspunkt zum AC3-Transformationspunkt durchgeführt wird.
11. Verfahren zur Herstellung eines hochfesten zweiphasigen verzinkten Stahlblechs mit
hervorragender Tiefziehfähigkeit nach Anspruch 10, dadurch gekennzeichnet, dass es ferner eine kontinuierliche Glühstufe zwischen der Kaltwalzstufe und der kontinuierlichen
Glühstufe in einem Temperaturbereich vom AC1-Transformationspunkt zum AC3-Transformationspunkt umfasst.
1. Tôle d'acier laminée à froid à deux phases à haute résistance ayant une excellente
aptitude à l'emboutissage, caractérisée en ce que la tôle d'acier a une composition comprenant C : 0,015 à 0,08 % en masse, Si : pas
plus de 2,0 % en masse, Mn : 0,5 à 3,0 % en masse, P : pas plus de 0,10 % en masse,
S : pas plus de 0,02 % en masse, Al : 0,005 à 0,20 % en masse, N : pas plus de 0,02
% en masse, V : 0,01 à 0,5 % en masse, éventuellement pas plus de 0,3 % en masse au
total d'un ou deux parmi Nb : plus de 0 % en masse mais pas plus de 0,3 % en masse,
et Ti : plus de 0 % en masse mais pas plus de 0,3 % en masse, et éventuellement pas
plus de 2,0 % en masse au total d'un ou deux parmi Cr et Mo, étant entendu que V,
Nb, Ti et C - s'ils sont présents - satisfont à une relation de 0,5xC/12 ≤ (V/51 +
2xNb/93 + 2xTi/48) ≤ 3xC/12, et le reste étant Fe et des impuretés inévitables, et
a une microstructure consistant en une phase de ferrite en tant que phase primaire
et une phase secondaire incluant une phase de martensite à une proportion de surface
non inférieure à 1 % par rapport à la totalité de la microstructure.
2. Tôle d'acier laminée à froid à deux phases à haute résistance ayant une excellente
aptitude à l'emboutissage selon la revendication 1, dans laquelle la tôle d'acier
comprend en outre pas plus de 0,3 % en masse au total d'un ou deux parmi Nb : plus
de 0 % en masse mais pas plus de 0,3 % en masse et Ti : plus de 0 % en masse mais
pas plus de 0,3 % en masse.
3. Tôle d'acier laminée à froid à deux phases à haute résistance ayant une excellente
aptitude à l'emboutissage selon la revendication 2, dans laquelle la tôle d'acier
comprend pas plus de 0,3 % en masse au total d'un ou deux parmi Nb : 0,001 à 0,3 %
en masse et Ti : 0,001 à 0,3 % en masse.
4. Tôle d'acier laminée à froid à deux phases à haute résistance ayant une excellente
aptitude à l'emboutissage selon la revendication 2, dans laquelle la tôle d'acier
comprend C : 0,03 à 0,08 % en masse, Si : 0,1 à 2,0 % en masse, Mn : 1,0 à 3,0 % en
masse, P : pas plus de 0,05 % en masse et S : pas plus de 0,01 % en masse, étant entendu
que V, Nb et Ti satisfont à une relation de 1,5 ≤ (2xNb/93 + 2xTi/48) / (V/51) ≤ 15.
5. Procédé de fabrication d'une tôle d'acier laminée à froid à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage, qui comprend un laminage à chaud d'une
brame d'acier ayant une composition comprenant : C : 0,015 à 0,08 % en masse, Si :
pas plus de 2,0 % en masse, Mn : 0,5 à 3,0 % en masse, P : pas plus de 0,10 % en masse,
S : pas plus de 0,02 % en masse, Al : 0,005 à 0,20 % en masse, N : pas plus de 0,02
% en masse, V : 0,01 à 0,5 % en masse, éventuellement pas plus de 0,3 % en masse au
total d'un ou deux parmi Nb : plus de 0 % en masse mais pas plus de 0,3 % en masse,
et Ti : plus de 0 % en masse mais pas plus de 0,3 % en masse, et éventuellement pas
plus de 2,0 % en masse au total d'un ou deux parmi Cr et Mo, étant entendu que V,
Nb, Ti et C - s'ils sont présents - satisfont à une relation de 0,5C/12 ≤ (V/51 +
2xNb/93 + 2xTi/48) ≤ 3xC/12, et le reste étant Fe et des impuretés inévitables, un
décapage, un laminage à froid, et puis une soumission à un recuit continu à gamme
de températures depuis un point de transformation AC1 jusqu'à un point de transformation AC3.
6. Procédé de fabrication d'une tôle d'acier laminée à froid à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage selon la revendication 5, dans lequel
la brame d'acier comprend en outre pas plus de 0,3 % en masse au total d'un ou deux
parmi Nb : plus de 0 % en masse mais pas plus de 0,3 % en masse et Ti : plus de 0
% en masse mais pas plus de 0,3 % en masse.
7. Procédé de fabrication d'une tôle d'acier laminée à froid à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage selon la revendication 6, dans lequel
la brame d'acier comprend pas plus de 0,3 % en masse au total d'un ou deux parmi Nb
: 0,001 à 0,3 % en masse et Ti : 0,001 à 0,3 % en masse.
8. Procédé de fabrication d'une tôle d'acier laminée à froid à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage selon la revendication 6, dans lequel
la brame d'acier comprend C : 0,03 à 0,08 % en masse, Si : 0,1 à 2,0 % en masse, Mn
: 1,0 à 3,0 % en masse, P : pas plus de 0,05 % en masse et S : pas plus de 0,01 %
en masse, étant entendu que V, Nb et Ti satisfont à une relation de 1,5 ≤ (2xNb/93
+ 2xTi/48) / (V/51) ≤ 15.
9. Tôle d'acier galvanisée à deux phases à haute résistance ayant une excellente aptitude
à l'emboutissage, comprenant un revêtement galvanisé sur la tôle d'acier telle que
revendiquée dans l'une quelconque des revendications 1 à 4.
10. Procédé de fabrication d'une tôle d'acier galvanisée à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage, caractérisé par le fait de soumettre à une galvanisation après le recuit continu à une gamme de températures
depuis un point de transformation AC1 jusqu'à un point de transformation AC3 dans le procédé dans l'une quelconque des revendications 5 à 8.
11. Procédé de fabrication d'une tôle d'acier galvanisée à deux phases à haute résistance
ayant une excellente aptitude à l'emboutissage selon la revendication 10, caractérisé en ce qu'il comprend en outre une étape de recuit continu entre l'étape de laminage à froid
et l'étape de recuit continu à une gamme de températures depuis un point de transformation
AC1 jusqu'à un point de transformation AC3.