1
[Name of Document] DESCRIPTION
[Title of the Invention] DUAL PHASE STEEL SHEET AND
MANUFACTURING METHOD THEREOF
[Technical Field]
5 [0001] The present invention relates to a dual phase steel sheet composed
of ferrite and low-temperature transformation products and a manufacturing
method thereof. This application is based upon and claims the benefit of
priority of the prior Japanese Patent Application No. 2012-212783, filed on
September 26, 2012, the entire contents of which are incorporated herein by
10 reference.
[Background Art]
[0002] In recent years, there has been promoted weight reduction of
various parts constituting an automobile in order to improve fuel consumption
of an automobile. Weight reduction means differ depending on each
15 required performance of the parts, and for example, for a framework part,
thickness thinning achieved by increasing strength of a steel sheet is
performed, and for a panel part, application of a light metal such as an Al
alloy to a steel sheet and the like are performed. However, when compared
to steel, the light metal such as an Al alloy is expensive, so that it is mainly
20 applied to luxury automobiles in the real world.
[0003] On the other hand, a demand for automobiles is being shifted to
emerging countries from developed countries, and from now on, it is expected
that weight reduction and price reduction are both achieved. For any parts, it
becomes necessary to achieve, of steel, strength increase and weight reduction
25 achieved by thickness thinning.
[0004] Aluminum casting and forgings have been advantageous to wheels
2
for passenger vehicles in terms of design. However, even though steel
pressed products are used recently as the wheels for passenger vehicles, by
devising materials and methods, products having the design equivalent to that
of an aluminum wheel are appearing.
5 [0005] Particularly, in addition to excellent fatigue endurance and
corrosion resistance that have been required so far in a wheel disc seen by an
end user, the design and beautifulness equivalent to those of an aluminum
wheel are also required in a steel wheel. Similarly, also in a steel sheet for
wheel disc, workability improvement for improving design as a part and
10 surface property improvement for securing beautifulness are required, in
addition to the strength increase that achieves thickness thinning, and the
fatigue endurance and the corrosion resistance that have been required so far.
[0006] As properties that have been required so far in the steel sheet for
wheel disc, bulging workability, drawability, and fatigue endurance have been
15 regarded as important in particular. This is because working of a hat portion
is challenging among forming steps of the wheel disc and the fatigue
endurance is managed by the strictest standard among member properties of
the wheel.
[0007] At present, in order to emphasize the fatigue endurance of a
20 member as a high-strength hot-rolled steel sheet for wheel disc,
ferrite-martensite dual phase steel sheets of 590 MPa grade excellent in
fatigue property (what is called Dual Phase steel) have been used. However,
the strength level required in these steel sheets is increased to the 780 MPa
grade from the 590 MPa grade and the strength tends to further increase.
25 [0008] In Non-Patent Document 1, there has been disclosed a method of
securing uniform elongation even with the same strength by turning a
3
microstructure of a steel sheet into a composite-structure such as a Dual Phase
steel composed of ferrite and martensite (to be described as DP steel,
hereinafter).
[0009] On the other hand, the DP steel has been known that local
5 deformability typified by bending forming, hole expansion, and burring is low.
This is because a strength difference between ferrite and martensite is large,
so that large strain and stress concentration occur in ferrite near martensite
with formation and a crack occurs.
[0010] Based on this finding, a high-strength steel sheet whose hole
10 expansion ratio is increased by decreasing the strength difference between
structures has been developed. In Patent Document 1, there has been
proposed a steel sheet in which strength is secured by applying bainite or
bainitic ferrite as its main phase to largely improve hole expandability. The
steel is designed to be composed of a single structure, and thereby the strain
15 and stress concentration described above are prevented from occurring and a
high hole expansion ratio can be obtained.
[0011] However, the steel is designed to be composed of a single
structure of bainite or bainitic ferrite, and thereby elongation deteriorates
greatly and the achievement of elongation and hole expandability cannot be
20 attained.
[0012] Further, in recent years, there have been proposed high-strength steel
sheets in which ferrite excellent in elongation is used as a structure of a single
structure steel and a strength increase is achieved by using carbide of Ti, Mo,
or the like (for example, Patent Documents 2 to 4).
25 [0013] However, the steel sheet proposed in Patent Document 2 contains
a large amount of Mo. The steel sheet proposed in Patent Document 3
4
contains a large amount of V. Further, the steel sheet proposed in Patent
Document 4 needs to be cooled in the middle of rolling for making ciystal
grains fine. Therefore, there is a problem that the alloy cost and the
manufacturing cost increase. Further, even in this steel sheet, ferrite itself is
5 largely increased in strength, and thereby the elongation deteriorates. The
elongation of the single structure steel composed of bainite or bainitic ferrite
is excellent, but the elongation-hole expandability balance is not necessarily
sufficient.
[0014] Further, in Patent Document 5, there has been proposed a dual
10 phase steel sheet in which in a DP steel, bainite is used in place of martensite
and a strength difference between structures of ferrite and bainite is decreased,
to thereby increase hole expandability.
[0015] However, as a result that an area ratio of the bainite structure was
increased in order to secure strength, the elongation deteriorated and the
15 elongation-hole expandability balance was not sufficient.
[0016] Further, in Patent Documents 7 to 9, there have been also
proposed steel sheets in which ferrite in a DP steel is
precipitation-strengthened and thereby a strength difference between ferrite
and hard structure is decreased.
20 [0017] However, in this technique, Mo is an essential element to cause a
problem that the manufacturing cost increases. Further, even though ferrite
is precipitation-strengthened, the strength difference between ferrite and
martensite being a hard structure is large, resulting in that a high hole
expandability improving effect is not obtained.
25 [0018] On the other hand, in order to turn a microstructure into a dual
phase of ferrite and martensite, Si is often added to these DP steels for the
5
purpose of promoting ferrite transformation. However, when Si is contained,
a tiger stripe scale pattern called a red scale (Si scale) is generated on the
surface of the steel sheet, so that it is difficult to apply the DP steel to various
steel sheets used for highly-designed wheel discs required to have
5 beautifulness.
[0019] In Patent Document 10, there has been disclosed a technique
relating to a steel sheet capable of obtaining an excellent balance between
elongation and hole expandability by controlling a martensite fraction in a DP
steel to 3 to 10% in a steel sheet of 780 MPa grade or higher. However,
10 0.5% or more of Si is added, thereby making it difficult to avoid the Si scale
pattern, so that it is difficult to apply the technique to various steel sheets used
for highly-designed wheel discs required to have beautifulness.
[0020] With regard to this problem, there has been disclosed a technique
of a high-tensile hot-rolled steel sheet capable of suppressing occurrence of
15 red scales by suppressing the added amount of Si to 0.3% or less and further
obtaining high strength and excellent stretch flangeability by adding Mo and
making precipitates fine (for example, Patent Documents 11 and 12).
[0021] However, in steel sheets having had the above-described
technique disclosed in Patent Documents 11 and 12 applied thereto, the added
20 amount of Si is about 0.3%) or less, but it is difficult to sufficiently suppress
occurrence of red scales, and further adding 0.07% or more of Mo being an
expensive alloy element is essential, so that there is a problem that the
manufacturing cost is high.
[0022] Further, in Patent Document 13, there has been disclosed a
25 technique of avoiding occurrence of red scales by defining the upper limit of
the content of Si. However, there is no technical disclosure on notch fatigue
6
property.
[0023] Further, in Patent Document 14, there has been disclosed a
technique of improving a low cycle fatigue property by adding Al. However,
there is no technical disclosure on notch fatigue property being a fatigue
5 property under stress concentration.
[Prior Art Document]
[Patent Document]
[0024] Patent Document 1: Japanese Laid-open Patent Publication No.
2003-193190
10 Patent Document 2: Japanese Laid-open Patent Publication No.
2003-089848
Patent Document 3: Japanese Laid-open Patent Publication No.
2007-063668
Patent Document 4: Japanese Laid-open Patent Publication No.
15 2004-143518
Patent Document 5: Japanese Laid-open Patent Publication No.
2004-204326
Patent Document 6: Japanese Laid-open Patent Publication No.
2007-302918
20 Patent Document 7: Japanese Laid-open Patent Publication No.
2003-321737
Patent Document 8: Japanese Laid-open Patent Publication No.
2003-321738
Patent Document 9: Japanese Laid-open Patent Publication No.
25 2003-321739
Patent Document 10: Japanese Laid-open Patent Publication No.
7
2011-184788
Patent Document 11: Japanese Laid-open Patent Publication No.
2002-322540
Patent Document 12: Japanese Laid-open Patent Publication No.
5 2002-322541
Patent Document 13: Japanese Patent Publication No. 2007-082567
Patent Document 14: Japanese Laid-open Patent Publication No.
2010-150581
[Non-Patent Document]
10 [0025] Non-Patent Document 1: O. Matsumura et al, Trans.
ISIJ(1987)vol. 27, p. 570
[Disclosure of the Invention]
[Problems to Be Solved by the Invention]
[0026] The present invention has an object to provide a high-burring
15 workability high-strength dual phase steel sheet having a tensile strength of
540 MPa or higher and having excellent surface property and notch fatigue
property and a manufacturing method thereof.
[Means for Solving the Problems]
[0027] The present inventors repeated earnest examinations on the
20 relationship between a structural constitution of a dual phase steel having a
high ductility as well as having a high strength and uniform elongation,
burring workability, and a notch fatigue property based on the premise of a
steel component not containing Si for the purpose of avoiding a Si scale
pattern. As a result, they found a method of bringing the uniform elongation,
25 the burring workability, and the notch fatigue property into balance on a high
level by controlling a steel component, a dispersion state, shape, size and
8
nanohardness of a low-temperature transformation product being a second
phase. That is, as a substitute for Si, Al was appropriately added to avoid a
Si scale pattern, and making a structure composite in which polygonal ferrite
is set as a main phase and a low-temperature transformation product is set as a
5 second phase was promoted. Further, they learned optimum ranges of a
fraction, a size, and the like of the low-temperature transformation product
that could achieve the elongation, the burring workability, and the notch
fatigue property. Further, they clarified that by devising not only the steel
component but also a hot rolling method, these optimum ranges can be
10 obtained with repeatability. The present invention has been made based on
such findings, and the gist thereof is as follows.
[0028] [1]
A dual phase steel sheet contains:
in mass%,
15 C: 0.01 to 0.1%;
Mn: 0.2 to 3%;
Al: 0.04 to 1.5%;
Ti: 0.015 to 0.2%;
Si: 0 to 0.5%;
20 Nb:0to0.06%;
Cu:0tol.2%;
Ni:0to0.6%;
Mo: 0 to 1%;
V: 0 to 0.2%;
25 Cr: 0 to 2%;
W: 0 to 0.5%;
9
Mg:0to0.01%;
Ca:0to0.01%;
REM: 0 to 0.1%;
B:0 to 0.002%;
5 P: 0.01% or less;
S: 0.005% or less;
N: 0.01% or less,
in which [Ti] - 48/14 x [N] - 48/32 x [S] ^ 0% is satisfied and when Ex.C
(%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]} is set,
10 0.001 ^ Ex.C (%)/fsd (%) £ 0.01 is satisfied, and
a balance being composed of Fe and impurities, in which
at the position of 1/4 thickness of a sheet thickness, a microstructure is
a dual phase with its main phase composed of polygonal ferrite
precipitation-strengthened by carbide of Ti and its second phase composed of
15 1 to 10%i in area fraction (fsd (%)) of low-temperature transformation
products dispersed plurally, and
an average crystal diameter of the low-temperature transformation
product is 3 to 15 urn and an average value of a distance of closest approach
between the low-temperature transformation products is 10 to 20 um.
20 [0029] [2] The dual phase steel sheet according to claim [1], contains:
in mass%,
Si: 0.02% to 0.5%.
[0030] [3] The dual phase steel sheet according to [1] or [2], contains:
one or two or more of
25 in mass%,
Nb: 0.005 to 0.06%;
10
Cu: 0.02 to 1.2%;
Ni: 0.01 to 0.6%;
Mo: 0.01 to 1%;
V: 0.01 to 0.2%;
5 CrO.Ol to 2%; and
W: 0.01 to 0.5%.
[0031] [4] The dual phase steel sheet according to any one of [1] to [3],
contains:
one or two or more of
10 in mass%,
Mg: 0.0005 to 0.01%;
Ca: 0.0005 to 0.01%; and
REM: 0.0005 to 0.1%.
[0032] [5] The dual phase steel sheet according to any one of claims [1]
15 to [4], contains:
in mass%,
B: 0.0002 to 0.002%.
[0033] [6] The dual phase steel sheet according to any one of [1] to [5], in
which
20 galvanizing is performed on its surface.
[0034] [7] A manufacturing method of a dual phase steel sheet includes:
on a slab containing:
in mass%,
C: 0.01 to 0.1%;
25 Mn: 0.2 to 3%;
Al: 0.04 to 1.5%;
11
Ti: 0.015 to 0.2% or less;
Si: 0 to 0.5%;
Nb: 0 to 0.06%;
Cu:0to 1.2%;
5 Ni:0to0.6%;
Mo:0tol%;
V: 0 to 0.2%;
Cr: 0 to 2%;
W:0to0.5%;
10 Mg:0to0.01%;
Ca:0to0.01%;
REM: 0 to 0.1%;
B:0 to 0.002%;
P: 0.01% or less;
15 S: 0.005% or less;
N: 0.01% or less,
in which [Ti] - 48/14 x [N] - 48/32 x [S] ^ 0% is satisfied and when Ex.C
(%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]} is set,
0.001 ^ Ex.C (%)/fsd (%) £ 0.01 is satisfied, and
20 a balance being composed of Fe and impurities, performing heating to a
temperature SRTmj„ (°C) or higher, which is defined by Expression (1) below,
and then in hot rolling, performing rough rolling at a reduction ratio of 20%o
or more in a temperature zone of not lower than 1050°C nor higher than
1150°C for at least one pass, and then starting finish rolling within 150
25 seconds in a temperature zone of 1000°C or higher and lower than 1080°C,
and completing finish rolling with the total reduction ratio for plural passes of
12
not less than 75% nor more than 95% in a temperature zone of not lower than
an Ar3 transformation point temperature + 50°C nor higher than 1000°C; and
within 3 seconds, performing cooling down to lower than the Ar3
transformation point temperature at an average cooling rate of 15°C/sec or
5 more, and next performing cooling down to a temperature zone of higher than
600°C at an average cooling rate of 10°C/sec or less for a time period of 1
second or longer and shorter than 100 seconds, and next performing cooling
down to a temperature zone of 350°C or lower at a cooling rate of 15°C/sec or
more, and performing coiling.
10 SRTmin - 10780/(5.13 - logflTi] x [C])} - 273 - Expression (1)
[0035] [8] The manufacturing method of the dual phase steel sheet
according to [7], further includes:
in the hot rolling, performing rough rolling at a reduction ratio of 20%
or more in a temperature zone of not lower than 1050°C nor higher than
15 1150°C for plural passes, in which
the total reduction ratio of the rough rolling is not less than 60% nor more
than 90%.
[0036] [9] The manufacturing method of the dual phase steel sheet
according to [7] or [8], further includes:
20 performing cooling down to a temperature zone of 100°C or lower and
performing coiling.
[0037] [10] The manufacturing method of the dual phase steel
sheet according to any one of [7] to [9], in which
in the performing the cooling down to the temperature zone of higher than
25 600°C at an average cooling rate of 10°C/sec or less for a time period of 1
second or longer and shorter than 100 seconds, when a total cumulative
13
diffusion length Ltotai of Ti in ferrite is expressed by Expression (3) below by
adding up a diffusion length L of Ti in ferrite expressed by Expression (2)
below for a very short time At/sec from a cooling completing temperature to
coiling, 0.15 ^ L[otai ^ 0.5 is satisfied.
5 L = VD(T + 273)t - Expression (2)
Ltotai = EV(D(T + 273)At) - Expression (3)
Here, D(T + 273) is a volume diffusion coefficient at T°C. t is a
diffusion time period.
D(T) is expressed by Expression (4) below using a diffusion coefficient DO of
10 Ti, an activation energy Q, and a gas constant R.
D(T) = DO x Exp(-Q/R-(T + 273)) - Expression (4)
[0038] [11] The manufacturing method of the dual phase steel
sheet according to any one of [7] to [10], in which
in the performing the cooling down to the temperature zone of higher than
15 600°C at an average cooling rate of 10°C/sec or less for a time period of 1
second or longer and shorter than 100 seconds, a steel sheet is immersed in a
galvanizing bath to galvanize its surface.
[0039] [12] The manufacturing method of the dual phase steel
sheet according to [11], further includes:
20 on a galvanized dual phase steel sheet, performing an alloying treatment in a
temperature range of 450 to 600°C.
[Effect of the Invention]
[0040] According to the present invention, it is possible to obtain a
high-strength dual phase steel sheet excellent in uniform elongation, burring
25 workability, and notch fatigue property, and further excellent also in surface
property as well as having a tensile strength of 540 MPa or higher, and
14
industrial contribution is extremely significant.
[Brief Description of the Drawings]
[0041] [FIG. 1] FIG. 1 is a view showing a notched fatigue test piece.
[Mode for Carrying out the Invention]
5 [0042] A dual phase steel sheet is a steel sheet in which hard
low-temperature transformation products typified by martensite are dispersed
in soft ferrite, and achieves high uniform elongation as well as being high in
strength. However, at the time of deformation, strain and stress
concentration caused by a strength difference between ferrite and martensite
10 occur, and voids to cause ductile fracture are likely to be generated to grow,
so that it is general that local deformabihty relating to burring workability is
quite low.
[0043] On the other hand, with regard to a notch fatigue property to
evaluate a fatigue property under stress concentration, it is known that most of
15 a fracture life is derived from propagation of a fatigue crack. In the dual
phase steel in which hard low-temperature transformation products typified
by martensite are dispersed in soft ferrite, it is conceivable that when a fatigue
crack propagates through soft ferrite, the hard low-temperature transformation
product becomes an obstacle to the fatigue crack propagation, propagation
20 speed decreases, and the notch fatigue property improves.
[0044] However, detailed examinations on a fraction, size, and the like of
the low-temperature transformation product in the dual phase steel sheet,
generation and growth behavior of voids causing the ductile fracture, and the
propagation speed of a fatigue crack are not conducted. The optimal
25 microstructure capable of achieving improvement in the local deformabihty
relating to the burring workability of the dual phase steel sheet and decrease
15
in the propagation speed of a fatigue crack is not necessarily definite.
[0045] Further, components and a manufacturing method of a steel sheet
capable of satisfying all of avoidance of a Si scale pattern relating to a surface
property of a steel sheet for the puipose of achieving the design and
5 beautifulness equivalent to those of an aluminum wheel with a steel wheel,
security of post-coating corrosion resistance, burring workability, and notch
fatigue property, are not necessarily definite.
[0046] Thus, the present inventors repeated earnest examinations on the
relationship between a structural constitution of a dual phase steel having a
10 high ductility as well as having a high strength and uniform elongation,
burring workability, and a notch fatigue property based on the premise of a
steel component not containing Si for the purpose of avoiding a Si scale
pattern. As a result, they found a method of bringing the uniform elongation,
the burring workability, and the notch fatigue property into balance on a high
15 level by controlling the steel component, the dispersion state, shape, size and
nanohardness of the low-temperature transformation product being a second
phase.
[0047] Concretely, the content of Si was controlled to 0.5% or less, to
thereby avoid the Si scale pattern. Further, in order to bring the area fraction
20 (fsd (%)), size and the like of the low-temperature transformation product into
appropriate ranges, the amount of Ex.C was controlled in a range satisfying
0.001 < Ex.C (%)/fsd (%) ^ 0.01 (being Ex.C (%) = [C] - 12/48 x {[Ti] +
48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]}, here). Further, at the position of
1/4 thickness of a sheet thickness, a microstructure was set to a dual phase
25 with its main phase composed of polygonal ferrite precipitation-strengthened
by carbide of Ti and its second phase composed of 1 to 10% in area fraction
16
(fsd (%)) of low-temperature transformation products dispersed plurally.
Then, an average crystal diameter of the aforesaid low-temperature
transformation product was set to 3 to 15 |im, and an average value of a
distance of closest approach between the low-temperature transformation
5 products was set to 10 to 20 urn. As a result, they made clear that it is
possible to bring the uniform elongation, the burring workability, and the
notch fatigue property into balance on a high level.
[0048] As a test method by which the difference of burring workability
appears clearly, a hole expanding test is proposed. A hole expansion value
10 obtained by this test is widely used as an index to evaluate the local
deformability relating to burring workability. Occurrence and progress of a
crack in hole expanding are caused by ductile fracture with generation,
growth, and connection of voids set as elementary steps. In a structure
having a large strength difference as is the dual phase steel sheet, high strain
15 and stress concentration occur due to hard low-temperature transformation
products, so that voids occur to grow easily and the hole expansion value is
low.
[0049] However, when the relationship between the structure and the
generation and growth behavior of voids and the relationship between them
20 and the hole expandability were examined in detail, it became clear that
depending on the dispersion state of low-temperature transformation product
being a hard second phase, the generation, growth, and connection of voids
are sometimes delayed, to thus make it possible to obtain an excellent hole
expansion value.
25 [0050] Concretely, when of the low-temperature transformation products
dispersed in an island shape, the area fraction fsd is 10% or less, the average
17
ciystal diameter is 15 jim or less, and the average value of the distance of
closest approach between the low-temperature transformation products is 20
um or less, the generation, growth, and connection of voids are delayed, to
thus make it possible to obtain an excellent hole expansion value.
5 [0051] This is because when the low-temperature transformation products
are made small and the number per unit volume is decreased, the
low-temperature transformation products being occurrence sites of voids
themselves or vicinities of boundaries between ferrite and the
low-temperature transformation products are decreased and respective
10 intervals between the low-temperature transformation products are increased,
and thereby voids are not easily connected and the growth of voids is
suppressed. Further, hardness of the low-temperature transformation
product is limited to a certain range, and thereby local occurrence of voids
being an initial stage of deformation can be avoided and non-uniform growth
15 of voids is suppressed.
[0052] On the other hand, the notch fatigue property can be improved by
dispersing the hard low-temperature transformation product and decreasing
the propagation speed of a fatigue crack. In the case of the dual phase steel,
it is known that the propagation speed of a fatigue crack changes depending
20 on the dispersion state of the low-temperature transformation product being a
hard second phase, and by optimizing the dispersion state, the effect is
exhibited.
[0053] Concretely, as long as of the low-temperature transformation
products dispersed in an island shape, the area fraction fsd is 1% or more, the
25 average crystal diameter is 3 um or more, and the average value of the
distance of closest approach between the low-temperature transformation
18
products is 10 u,m or more, a fatigue crack to go through soft ferrite stays at or
bypasses the low-temperature transformation product being a hard second
phase, and thereby the propagation speed of the fatigue crack decreases and
notch fatigue strength improves.
5 [0054] Further, as long as the low-temperature transformation products
being a second phase have the average crystal diameter of 3 to 15 jam and
have the average value of the distance of closest approach therebetween of 10
to 20 urn, and are in a state of being dispersed in an island shape in an area
fraction of 1 to 10%, excellent uniform elongation that the dual phase steel
10 exhibits can be obtained.
[0055] In the foregoing, the characteristics of the present invention have
been explained in principle, and there will be next explained requirements
defining the present invention and preferable requirements sequentially.
First, components of the present invention will be explained in detail.
15 Incidentally, with regard to the component, % means mass%.
[0056] C: 0.01 to 0.1%
C is one of important elements in the present invention. C not only
forms low-temperature transformation products to contribute to strength by
structure strengthening, but also forms precipitates with Ti to contribute to
20 strength by precipitation strengthening. However, when C is less than
0.01%, these effects for securing the strength of 540 MPa or higher cannot be
obtained. When greater than 0.1% of C is contained, an area ratio of the
low-temperature transformation product being a hard second phase is
increased and the hole expandability decreases. Thus, the content of C is set
25 to 0.01% to 0.1%.
[0057] Further, as long as 0.001 ^ Ex.C (%)/fsd (%) ^ 0.01 (Ex.C
19
(%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]}) is
satisfied on the condition that the area fraction of the second phase is set to
fsd (%), the dispersion state, hardness, and the like of the low-temperature
transformation product being a hard second phase are optimized, the
5 generation, growth, and connection of voids are delayed, an excellent hole
expansion value can be obtained, and the tip of a fatigue crack stays or makes
a detour, and thereby the propagation speed of the fatigue crack decreases and
excellent notch fatigue strength can be obtained. Incidentally, in the
expression expressing Ex.C (%), [C] is the content of C (mass%), [Ti] is the
10 content of Ti (mass%), [Nb] is the content of Nb (mass%), [N] is the content
of N (mass%), and [S] is the content of S (mass%).
[0058] Mn: 0.2 to 3%
Mn is not only an element involved in strengthening of ferrite, but
also an element expanding an austenite region temperature to a low
15 temperature side to expand a two-phase region temperature zone of ferrite and
austenite with an increase in its content. In order to obtain the dual phase
steel of the present invention, it is necessaiy to promote two-phase separation
of ferrite and austenite during cooling after finish rolling. In order to obtain
the effect, 0.2% or more of Mn needs to be contained. On the other hand,
20 when Mn is contained in excess of 3%, slab cracking significantly occurs
during casting, so that the content is set to 3% or less.
[0059] Further, when greater than 2.5% of Mn is contained, hardenability
increases too much, resulting in that an intended microstructure cannot be
obtained by an ordinary method. In order to obtain the intended
25 microstructure, air-cooling and holding for a long time is required for
precipitating ferrite during cooling after finish rolling, and productivity
20
decreases, so that the content is desirably 2.5% or less. It is further desirably
2.2% or less. Further, when elements other than Mn are not added
sufficiently for the purpose of suppressing occurrence of hot cracking caused
by S, the amount of Mn that makes the content of Mn ([Mn]) and the content
5 of S ([S]) satisfy [Mn]/[S] ^ 20 in mass % is desirably contained.
[0060] Al: 0.04 to 1.5%
Al is involved in generation of ferrite similarly to Si to be one of
important elements in the present invention as well as being a deoxidizing
element. Al is also an element that with an increase in its content, expands a
10 ferrite region temperature to a high-temperature side to expand a two-phase
region temperature zone of ferrite and austenite, so that it is actively
contained as a substitute for Si in the present invention. In order to obtain
the effect, 0.04% or more of Al needs to be contained, but when it is
contained in excess of 1.5%, the ferrite region temperature is expanded to the
15 high-temperature side too much to thereby make it difficult to complete finish
rolling in an austenite region, and worked ferrite remains in a product sheet
and ductility deteriorates. Thus, the content of Al is set to not less than
0.04% nor more than 1.5%. Further, when greater than 1% of Al is
contained, there is caused a risk that non-metal inclusions such as alumina are
20 increased to deteriorate local ductility, so that it is desirably 1% or less.
[0061] Ti: 0.015 to 0.2%
Ti is one of the most important elements in the present invention.
Simultaneously with ferrite transformation progressing during cooling after
completion of hot rolling, the rest of Ti after having been precipitated as TiN
25 in an austenite region during hot rolling finely precipitates as carbide such as
TiC to precipitation strengthen ferrite grains of the dual phase steel of the
21
present invention, and thereby strength is improved. In order to obtain this
effect, Ti that is 0.015% or more and satisfies [Ti] - 48/14 x [N] - 48/32 x [S]
^ 0% needs to be contained.
[0062] On the other hand, even when greater than 0.2% of Ti is contained,
5 these effects are saturated. Further, 0.001 ^ Ex.C (%)/fsd (%) ^ 0.01
(Ex.C (%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]})
is set on the condition that the area fraction of the second phase is set to fsd
(%), and thereby the dispersion state, hardness, and the like of the
low-temperature transformation product being a hard second phase are
10 optimized, the generation, growth, and connection of voids are delayed, and
an excellent hole expansion value can be obtained. Further, the tip of a
fatigue crack stays at the low-temperature transformation product or bypasses
the low-temperature transformation product, and thereby the propagation
speed of the fatigue crack decreases and excellent notch fatigue strength can
15 be obtained. Further, when greater than 0.15% of Ti is contained, there is
caused a risk that a tundish nozzle is likely to be clogged at the time of
casting, so that it is desirably 0.15%o or less.
[0063] The steel used for the steel sheet of the present invention contains
the above elements as essential components, and further may also contain Si,
20 Nb, Cu, Ni, Mo, V, Cr, W, Mg, Ca, REM, and B according to need. These
respective elements will be described below.
[0064] Si: 0 to 0.5%
In the present invention, Si is not essential. Si is involved in
generation of ferrite as well as being a deoxidizing element, and is an element
25 that with an increase in its content, expands a ferrite region temperature to a
high-temperature side to expand a two-phase region temperature zone of
22
ferrite and austenite. In order to obtain the dual phase steel of the present
invention, Si is desirably contained originally. However, Si noticeably
generates a tiger stripe Si scale pattern on the surface of the steel sheet to
deteriorate surface property significantly. Further, there is sometimes a case
5 that it extremely decreases productivity of a scale removing step (pickling and
the like) on a precise adjustment line.
[0065] When greater than 0.07% of Si is contained, the Si scale pattern
begins to be found here and there on the surface of the steel sheet. When it
is greater than 0.5%, the surface property deteriorates significantly and the
10 productivity of a pickling step deteriorates extremely. Even though any
scale removing method is performed, a conversion treatment property
deteriorates and post-coating corrosion resistance decreases. Thus, the
content of Si is set to 0.5% or less.
[0066] On the other hand, Si is an element having an effect of
15 suppressing occurrence of scale-based defects such as scales and spindle
scales, and when 0.02% or more is contained, the effect can be obtained.
However, even though Si is contained in excess of 0.1%, the effect is
saturated, and furthermore the conversion treatment property deteriorates and
the post-coating corrosion resistance decreases. Thus, when Si is contained,
20 the content of Si is set to not less than 0.02% nor more than 0.5%, and is
desirably 0.1% or less. Further, in order to make the Si scale patterns zero,
the content of Si is desirably 0.07%o or less. However, the scale-based
defects such as scales and spindle scales vary in grade depending on needs,
and Si may also be less than 0.02%. A steel component not containing Si is
25 also in the range of the present invention.
[0067] • One or two or more of Nb, Cu, Ni, Mo, V, Cr, and W
23
In the present invention, Nb, Cu, Ni, Mo, V, Cr, and W are not
essential. Nb, Cu, Ni, Mo, V, Cr, and W are elements effective for
improving the strength of the steel sheet by precipitation-strengthening or
solid-solution strengthening. Therefore, one or two or more of Nb, Cu, Ni,
5 Mo, V, Cr, and W are contained according to need. When the content of Nb
is less than 0.005%, the content of Cu is less than 0.02%, the content of Ni is
less than 0.01%, the content of Mo is less than 0.01%, the content of V is less
than 0.01%, the content of Cr is less than 0.01%, and the content of W is less
than 0.01%, the above-described effect cannot be obtained sufficiently.
10 Further, even when greater than 0.06% of the content of Nb, greater than
1.2% of the content of Cu, greater than 0.6% of the content of Ni, greater than
1% of the content of Mo, greater than 0.2% of the content of V, greater than
2% of the content of Cr, and greater than 0.5% of the content of W are each
added, the above-described effect is saturated and economic efficiency
15 decreases.
[0068] Thus, when these are contained according to need, the content of
Nb is desirably not less than 0.005% nor more than 0.06%, the content of Cu
is desirably not less than 0.02% nor more than 1.2%, the content of Ni is
desirably not less than 0.01% nor more than 0.6%, the content of Mo is
20 desirably not less than 0.01% nor more than 1%, the content of V is desirably
not less than 0.01%) nor more than 0.2%, the content of Cr is desirably not
less than 0.01%> nor more than 2%, and the content of W is desirably not less
than 0.01% nor more than 0.5%.
[0069] One or two or more of Mg, Ca, and REM
25 In the present invention, Mg, Ca, and REM are not essential. Mg, Ca,
and REM (rare-earth element) are elements that control form of a non-metal
24
inclusion to be a starting point of fracture and to cause deterioration of
workability and improve workability. Therefore, one or two or more of Mg,
Ca, and REM are contained according to need. Even when less than
0.0005% of each of Ca, REM, and Mg is contained, the above-described
5 effect is not exhibited. Further, even when the content of Mg is set to greater
than 0.01%, the content of Ca is set to greater than 0.01%, and the content of
REM is set to greater than 0.1%, the above-described effect is saturated and
economic efficiency decreases.
[0070] Thus, when these are contained according to need, the content of
10 Mg is desirably not less than 0.0005% nor more than 0.01%, the content of Ca
is desirably not less than 0.0005% nor more than 0.01%, and the content of
REM is desirably not less than 0.0005% nor more than 0.1%. Incidentally,
in the present invention, REM refers to an element of La and the lanthanide
series, is often added in misch metal, and contains elements of the series such
15 as La and Ce in a complex form. Metals La and Ce may also be contained.
[0071] B: 0.0002 to 0.002%
In the present invention, B is not essential. B has an effect of
increasing hardenability to increase a structural fraction of a low-temperature
transformation generating phase being a hard phase, to thus be contained
20 according to need. However, when B is less than 0.0002%, the effect cannot
be obtained, and even though B is contained in excess of 0.002%, the effect is
saturated. Therefore, the content of B is desirably not less than 0.0002% nor
more than 0.002%). On the other hand, B is an element that causes concern
of slab cracking in a cooling step after continuous casting, and from this point
25 of view, the content is desirably 0.0015% or less. That is, it is desirably not
less than 0.001% nor more than 0.0015%.
25
[0072] With regard to the steel component of a hot-rolled steel sheet of
the present invention, its balance other than the above-described elements is
Fe and impurities. As the impurities, one contained in a raw material of ore,
scrap, and the like and one contained in a manufacturing step can be
5 exemplified. It is allowable that respective impurity elements are contained
as necessaiy in a range where the operation and effect of the present invention
are not inhibited.
[0073] P: 0.01% or less
P is an impurity element, and when it exceeds 0.01%, segregation to
10 ciystai grain boundaries becomes noticeable, grain boundary embrittlement is
promoted, and local ductility deteriorates. Further, embrittlement of a
welded portion also becomes noticeable, so that the upper limit is set to
0.01% or less. The lower limit value of P is not defined in particular, but
setting it to less than 0.0001% is economically disadvantageous.
15 [0074] S: 0.005% or less
S is an impurity element, and adversely affects weldabiiity and
manufacturability during casting and manufacturability during hot rolling, so
that the upper limit is set to 0.005% or less. Further, when S is contained
excessively, coarse MnS is formed to decrease hole expandability, so that for
20 improvement in hole expandability, the content is preferably decreased. The
lower limit value of S is not defined in particular, but setting it to less than
0.0001% is disadvantageous economically, so that this value is preferably set
to the lower limit value.
[0075] N: 0.01% or less
25 N is an impurity element to be mixed inevitably at the time of refining
of steel, and is an element to form nitride combined with Ti, Nb, or the like.
26
When the content of N is greater than 0.01%, this nitride precipitates at
relatively high temperature, so that crystal grains are likely to become coarse,
and the coarse crystal grain might become a starting point of a burring crack.
Further, this nitride is preferably less in order to effectively use Nb and Ti as
5 will be described later. Thus, the upper limit of the content of N is set to
0.01%.
[0076] Incidentally, when the content of N is greater than 0.006% in
applying the present invention to a member in which aging deterioration
becomes a problem, the aging deterioration becomes severe, so that it is
10 desirably 0.006% or less. Further, when the present invention is applied to a
member based on the premise that it is allowed to stand at room temperature
for two weeks or longer after manufacture, to then be subjected to working,
the content of N is desirably 0.005% or less in view of aging deterioration
measures. Further, when it is considered that a member is allowed to stand
15 under a summer high-temperature environment or it is used under an
environment with export to regions located over the equator by ships, vessels,
and the like, the content of N is desirably less than 0.004%.
[0077] As the other impurities, 1% or less in total of Zr, Sn, Co, and Zn
may also be contained. However, Sn is desirably 0.05% or less because a
20 flaw might occur at the time of hot rolling.
[0078] Subsequently, the microstructure of the dual phase steel sheet of
the present invention will be explained in detail. The microstructure of the
dual phase steel sheet of the present invention is limited as follows.
[0079] At the position of 1/4 thickness of a sheet thickness, the
25 microstructure is a dual phase with its main phase composed of polygonal
ferrite precipitation-strengthened by carbide of Ti and its second phase
27
composed of 1 to 10% in area fraction (fsd (%)) of low-temperature
transformation products dispersed plurally. An average crystal diameter of
the aforesaid low-temperature transformation product is 3 to 15 urn. An
average value of a distance of closest approach between the low-temperature
5 transformation products is 10 to 20 um. Incidentally, the microstructure is
specified at the position of 1/4 thickness of the sheet thickness where average
characteristics appear.
[0080] Ferrite is the most important structure for securing uniform
elongation. In order to obtain the strength of 540 MPa grade or higher even
10 when the area fraction of the low-temperature transformation product being a
hard second phase is 10% or less, the ferrite structure needs to be
strengthened by precipitation strengthening. Further, in order to secure
elongation, it is important that the main phase of the microstructure is not
bainitic ferrite having a high dislocation density but polygonal ferrite having a
15 low dislocation density and having sufficient ductility. Thus, the main phase
of the steel of the present invention is set to polygonal ferrite
precipitation-strengthened by carbide of Ti. Incidentally, the carbide of Ti to
be said here is a compound having Ti and C contributing to precipitation
strengthening of the ferrite structure as its main component, and it is also
20 acceptable to contain, for example, N, V, Mo, and the like in addition to Ti
andC.
[0081] As long as the component is fixed, the average grain diameter and
the density (piece/cm3) of precipitates containing TiC are substantially
inversely-correlated. In order for an improved margin of the strength by
25 precipitation strengthening to become 100 MPa or higher in terms of tensile
strength, of the precipitates containing TiC, the average grain diameter needs
28
1 f t
to be 3 nm or less and the density needs to be 1 x 10 pieces/cm or more.
[0082] In the present invention, the low-temperature transformation
product being a hard second phase is mainly martensite or bainite (ctB) not
containing coarse carbide between laths. However, it is allowable to contain
5 less than 3% in total in area ratio of retained austenite (yr) and
Martensite-Austenite constituent (MA). Further, the martensite to be said in
the present invention is fresh martensite (M) when coiling is performed in a
temperature zone of 100°C or lower where a diffusion speed of carbon is
sufficiently slow. It is tempered martensite (tM) when a coiling temperature
10 is higher than 100°C and an Ms point (an Ms point of remaining austenite
obtained after ferrite transformation progresses during cooling after finish
rolling) or lower. The low-temperature transformation product in the latter
case is a structure mixed with tempered martensite and bainite.
[0083] The ratio of tempered martensite and bainite of this mixed
15 structure (low-temperature transformation product in the latter case) is
affected by the coiling temperature and the relative relationship between the
coiling temperature and the above-described Ms point temperature.
Incidentally, when the Ms point is lower than 350°C, most of the
low-temperature transformation product is bainite not containing coarse
20 carbide between laths that is transformed at higher than the Ms point and
350°C or lower. However, it is metallographically difficult to distinguish
tempered martensite and bainite to be said here, and in the present invention,
these are referred to as tempered martensite (tM).
[0084] The low-temperature transformation product needs to be dispersed
25 in an island shape at a corner, an edge, and a grain boundary of a ferrite grain.
This is because with regard to ductile fracture thought to be involved in
29
burring workability, in a mechanism in which voids occur and then grow to be
connected, the shape of the low-temperature transformation product itself
thought to be an occurrence site of a void is an island shape, and thereby
stress concentration is relaxed and the occurrence of voids causing fracture of
5 the low-temperature transformation product is suppressed.
[0085] Incidentally, the island shape indicates a state where
low-temperature transformation products are not arranged continuously in an
aligned manner, and further the individual shape of them is desirably a shape
close to a sphere with few stress concentration places. As long as the
10 average ciystai diameter of the low-temperature transformation product is 3 to
15 um and the average value of the distance of closest approach between the
low-temperature transformation products is 10 to 20 um, the low-temperature
transformation products each have an appropriate size and are appropriately
dispersed to be in an "island shape."
15 [0086] Further, the low-temperature transformation product being a hard
second phase is an important structure in terms of securing uniform
elongation. When the area fraction (fsd (%)) of the low-temperature
transformation products dispersed in an island shape becomes less than 1%, it
becomes difficult to secure 15% or more of uniform elongation at the 540
20 MPa grade, for example. Further, an effect of delaying the propagation of a
fatigue crack is lost. On the other hand, when it becomes greater than 10%,
the intervals between the low-temperature transformation products thought to
be occurrence sites of voids become short, voids are likely to be connected,
ductile fracture is likely to be caused, and the burring workability deteriorates.
25 Therefore, the area fraction (fsd (%)) of the low-temperature transformation
product in the microstructure is limited to 1 to 10%.
30
[0087] The average crystal diameter of the low-temperature
transformation product needs to be limited to 3 to 15 um in terms of circle
equivalent diameter. This is because when the average crystal diameter of
the low-temperature transformation product is less than 3 um, the effect that
5 the low-temperature transformation product becomes an obstacle to the
propagation of a fatigue crack to delay the propagation speed is lost, and
when it is greater than 15 urn, the shape becomes complex naturally, stress
concentration portions are generated, fracture of a coarse low-temperature
transformation product is caused early, and local ductile fracture caused by
10 occurrence of voids adversely affects burring workability. It is desirably 12
urn or less.
[0088] Further, the average value of the distance of closest approach
between the low-temperature transformation products needs to be limited to
10 to 20 um. When the average value of the distance of closest approach
15 between the low-temperature transformation products is less than 10 um, the
intervals between the low-temperature transformation products become short,
voids are likely to be connected, ductile fracture is likely to be caused, and the
burring workability deteriorates. On the other hand, when the average value
of the distance of closest approach between the low-temperature
20 transformation products is greater than 20 um, a fatigue crack selectively
propagates through soft polygonal ferrite, and the effect of delaying the
propagation of a fatigue crack is lost.
[0089] The average nanohardness of the low-temperature transformation
product is desirably 7 to 18 GPa. This is because when the average
25 nanohardness is less than 7 GPa, a hardness difference between the
low-temperature transformation product and a soft ferrite phase is decreased
31
and excellent uniform elongation being the characteristic of the dual phase
steel is not exhibited. On the other hand, when it is greater than 18 GPa, the
hardness difference between the low-temperature transformation product and
a soft ferrite phase is increased by contraries, and voids occur locally at the
5 initial stage of deformation, and thus ductile fracture is likely to develop and
local deformability decreases. Further, a nanohardness range becomes 1.2
GPa or less in terms of standard deviation, and thereby the local occurrence of
voids at the initial stage of deformation is suppressed.
[0090] Sequentially, there will be explained a manufacturing method of s
10 steel sheet of the present invention.
In the present invention, a manufacturing method of a steel billet
(slab) having the above-described components to be performed before a hot
rolling step is not limited in particular. That is, as a manufacturing method
of a steel billet (slab) having the above-described components, it may also be
15 set that subsequently to a melting step by a shaft furnace, a converter, an
electric furnace, or the like, component adjustment is variously performed so
as to obtain intended component contents in a secondary refining step, and
next a casting step is performed by normal continuous casting, casting by an
ingot method, or a method of thin slab casting or the like. Incidentally, scrap
20 may also be used for a raw material. Further, when a slab is obtained by
continuous casting, an intact high-temperature cast slab may be directly
transformed to hot rolling, or the slab may also be hot rolled after being
cooled down to room temperature to then be reheated in a heating furnace.
[0091] The slab obtained by the above-described manufacturing method
25 is heated in a heating furnace at a minimum slab reheating temperature (=
SRTmin) or higher, which is calculated based on Expression (1), in a slab
32
heating step before hot rolling.
SRTmill = 10780/(5.13 - logflTi] x [C])} - 273 - Expression (1)
When it is lower than this temperature, carbonitride of Ti is not
sufficiently melted in a parent material. In this case, it is not possible to
5 obtain an effect that strength is improved by using precipitation strengthening
obtained by fine precipitation of Ti as carbide during cooling after completion
of finish rolling or after coiling. Thus, the heating temperature in the slab
heating step is set to the minimum slab reheating temperature (= SRTmin) or
higher, which is calculated in Expression (1). Incidentally, when the heating
10 temperature is lower than 1100°C, operational efficiency is significantly
impaired in terms of a schedule, so that the heating temperature is desirably
1100°C or higher.
[0092] Further, a heating time in the slab heating step is not defined in
particular, but in order to sufficiently promote the melting of carbonitride of
15 Ti, after the temperature reaching the above-described heating temperature,
the slab is desirably held for 30 minutes or longer. Further, when the slab is
sufficiently uniformly heated in a thickness direction of the slab, it is
desirably held for 60 minutes or longer. On the other hand, in terms of a
decrease in yield caused by scale off, it is 240 minutes or shorter. However,
20 when the cast slab obtained after casting is directly transferred to be rolled in
a high temperature state, the above is not applied.
[0093] After the slab heating step, on the slab extracted from the heating
furnace, a rough rolling step of hot rolling is started with no waiting time in
particular, and a rough bar is obtained. In this rough rolling step, rough
25 rolling at a rolling ratio of at least 20% or more needs to be performed for at
least one pass in a temperature zone of not lower than 1050°C nor higher than
33
1150°C.
[0094] When a rough rolling completing temperature is lower than
1050°C, hot deformation resistance during the rough rolling increases,
resulting in that operation of the rough rolling might be damaged. When it is
5 higher than 1150°C, secondaiy scales to be generated during the rough rolling
grow too much, resulting in that descaling to be performed later and removing
scales in finish rolling might be difficult to be performed.
[0095] Further, unless the roiling at a rolling ratio of 20% or more is
performed in the rough rolling in the temperature zone, refining of crystal
10 grains using working and subsequent recrystallization of austenite, and
resolution of anisotropy caused by a solidified structure cannot be expected.
Thereby, transformation behavior after finish rolling is affected, the shape of
the low-temperature transformation product being a second phase in the
microstructure of the dual phase steel sheet changes to a film shape from an
15 island shape, and the burring workability deteriorates. Further, when the
cast slab obtained after casting is directly transferred to be rolled in a high
temperature state, a cast structure remains, and the shape change of the
low-temperature transformation product being a second phase to the film
shape might be noticeable.
20 [0096] The number of rolling passes in the rough rolling is preferably
plural passes, which is two passes or more. When plural passes are applied,
working and recrystallization in austenite are performed repeatedly and
average austenite grains before finish rolling are refined to 100 |im or less,
resulting in that the average grain diameter of the low-temperature
25 transformation product being a hard second phase is made 12 um or less
stably.
34
[0097] Further, the total reduction ratio in the rough rolling is preferably
60% or more. When the total reduction ratio is less than 60%, the
above-described effect of refining austenite grains cannot be obtained
sufficiently. However, even when the total reduction ratio in the rough
5 rolling is greater than 90%, the effect is saturated and further the number of
passes is increased to impede productivity, and a temperature decrease might
be caused. Further, due to the similar reason, the number of passes is
desirably 11 or less.
[0098] Finish rolling is performed after completion of the rough rolling.
10 The time period until start of finish rolling after completion of the rough
rolling is within 150 seconds.
[0099] When this time period is longer than 150 seconds, in the rough bar,
Ti in austenite precipitates as coarse carbide of TiC. As a result, the amount
of TiC to finely precipitate in ferrite at the time of austenite/ferrite
15 transformation during cooling to be performed later or at the time of
completion of ferrite transformation after coiling and to contribute to strength
by precipitation strengthening decreases and the strength decreases.
Furthermore, grain growth of austenite progresses and thereby the average
austenite grains before finish rolling become coarse to be greater than 100 |im,
20 resulting in that the average grain diameter of the low-temperature
transformation product being a hard second phase is sometimes made greater
than 15 urn.
[0100] On the other hand, the lower limit value of the time period until
start of finish rolling after completion of the rough rolling does not have to be
25 limited in particular. However, when it is shorter than 30 seconds, a finish
rolling start temperature does not decrease to lower than 1080°C unless a
35
special cooling device is used, and blisters to be a starting point of scales and
spindle scale defects occur between the surface of a base iron of the steel
sheet and scales before finish rolling and during passes, so that these scale
defects might be likely to be generated. Thus, it is desirably 30 seconds or
5 longer. .
[0101] A rolling start temperature of the finish rolling is set to 1000°C or
higher and lower than 1080°C.
[0102] When this temperature is lower than 1000°C, Ti precipitates in
austenite as coarse carbide of TiC by strain-induced precipitation during the
10 finish rolling. As a result, the amount of TiC to finely precipitate in ferrite at
the time of austenite/ferrite transformation during cooling to be performed
later or at the time of completion of ferrite transformation after coiling and to
contribute to strength by precipitation strengthening decreases and the
strength decreases.
15 [0103] On the other hand, when this temperature is higher than 1080°C,
blisters to be a starting point of scales and spindle scale defects occur between
the surface of a base iron of the steel sheet and scales before finish rolling and
during passes, so that these scale defects might be likely to be generated.
[0104] A finish rolling completing temperature is set to not lower than an
20 Ar3 transformation point temperature + 50°C nor higher than 1000°C.
[0105] The Ar3 transformation point temperature is simply expressed by,
for example, the following calculation expression in relation to the steel
components. That is, it is described by Expression (5) below.
Ar3 = 910 - 310 x [C] + 25 x {[Si] + 2 x [Al]} - 80 x [Mneq] -
25 Expression (5)
[0106] Here, when B is not added, [Mneq] is expressed by Expression (6)
36
below.
[Mneq] = [Mn] + [Cr] + [Cu] + [Mo] + [Ni]/2 + 10([Nb] - 0.02) -
Expression (6)
[0107] Further, when B is added, [Mneq] is expressed by Expression (7)
5 below.
[Mneq] - [Mn] + [Cr] + [Cu] + [Mo] + [Ni]/2 + 10([Nb] - 0.02) + 1 -
Expression (7)
[0108] Incidentally, [C] is the content of C (mass%), [Si] is the content of
Si (mass%), [Al] is the content of Al (mass%), [Cr] is the content of Cr
10 (mass%), [Cu] is the content of Cu (mass%), [Mo] is the content of Mo
(mass%), [Ni] is the content of Ni (mass%), and [Nb] is the content of Nb
(mass%).
[0109] When the finish rolling completing temperature is lower than the
At*3 transformation point temperature + 50°C, the low-temperature
15 transformation products in the microstmcture of the dual phase steel sheet are
brought into a dispersion state where they are continuously arranged in an
aligned manner. Furthermore, the average value of the distance of closest
approach between the low-temperature transformation products becomes less
than 10 urn, voids are likely to be connected, ductile fracture is likely to be
20 caused, and the burring workability deteriorates.
[0110] On the other hand, when it is higher than 1000°C, even when a
cooling pattern after rolling is controlled in any way, ferrite transformation
becomes insufficient and the area fraction of the low-temperature
transformation product in the microstmcture of a product sheet becomes
25 greater than 10%, and the burring workability deteriorates after all.
[0111] Further, the finish rolling is rolling with plural passes by a tandem
37
mill, and the total reduction ratio is not less than 75% nor more than 95%.
[0112] As long as the finish rolling is performed in a tandem mill
enabling rolling with plural passes, reduction is performed through plural
passes in the rolling, and thereby non-recrystallization by rolling and
5 recrystallization for an inter-pass time period until the coming pass are
repeated plural times. As a result, austenite grains are refined and the
average grain diameter of the low-temperature transformation product in the
microstructure of the dual phase steel sheet can be made 15 \xm or less.
However, when the total reduction ratio is less than 75%, austenite grains
10 cannot be refined sufficiently and the average grain diameter of the
low-temperature transformation product in the microstructure of the dual
phase steel sheet cannot be made 15 pm or less.
[0113] On the other hand, when it is greater than 95%, the effect is
saturated, and further an excessive load is applied to the rolling mill, so that it
15 is not preferable operationally.
[0114] Further, a reduction ratio in each pass is desirably 10% or more.
When the reduction ratio in each pass is less than 10% for three passes at the
rear stand of a finishing mill in particular and an average rolling ratio for three
passes is less than 10%, grain growth progresses significantly during the three
20 passes and after completion of the finish rolling, and there is a risk that the
average grain diameter of the low-temperature transformation product in the
microstructure of the dual phase steel sheet is no longer able to be made 12
urn or less.
[0115] Incidentally, in the present invention, a rolling speed is not limited
25 in particular. However, when the rolling speed at a finish final stand is less
than 400 mpm, the time period for each finish rolling pass is prolonged. As
38
a result, austenite grains grow to be coarse, and there is a risk that the average
grain diameter of the low-temperature transformation product in the
microstructure of a product sheet is no longer able to be made 15 urn or less
stably. Therefore, the rolling speed is desirably 400 mpm or more. Further,
5 when it is 650 mpm, the average grain diameter of the low-temperature
transformation product can be made 12 u,m or less stably, so that 650 mpm is
further desirable. Further, even if the upper limit is not limited in particular,
the effect of the present invention is achieved, but it is realistically 1800 mpm
or less due to facility restriction.
10 [0116] After completion of the finish rolling, in order to elaborate the
microstructure of a product, cooling optimized by control of a run-out-table is
performed.
First, the time period until start of the cooling after completion of the
finish rolling is within three seconds. When this time period until start of the
15 cooling is longer than three seconds, in austenite before being transformed,
precipitation of coarse and unaligned carbonitride of Ti progresses, the
precipitation amount of fine and aligned carbide of Ti to precipitate in ferrite
during cooling to be performed later decreases, and the strength might be
decreased. Further, austenite grains grow to be coarse, and there is a risk
20 that the average grain diameter of the low-temperature transformation product
in the microstructure of the product sheet is no longer able to be made 15 um
or less.
[0117] The lower limit value of the time period until start of this cooling
does not have to be limited in particular in the present invention, but when it
25 is shorter than 0.4 seconds, cooling is performed in a state where a lamellar
worked structure obtained by rolling remains, even in a product sheet,
39
low-temperature transformation products continuously arranged in an aligned
manner are obtained, and the burring workability might deteriorate.
[0118] As for the rate of a first-stage cooling step to be first performed
after completion of the rolling, an average cooling rate of 15°C/sec or more is
5 required. When this cooling rate is less than 15°C/sec, pearlite is formed
during cooling, and an intended microstructure might not be obtained.
Incidentally, even if the upper limit of the cooling rate in the first-stage
cooling step is not limited in particular, the effect of the present invention can
be obtained. However, when the cooling rate is greater than 150°C/sec,
10 controlling a cooling completing temperature is extremely difficult to make it
difficult to elaborate the microstructure, so that it is desirably set to 150°C/sec
or less.
[0119] A cooling stop temperature in the first-stage cooling step is lower
than the Ar3 transformation point temperature. When the cooling stop
15 temperature is the Ar3 transformation point temperature or higher, it is not
possible to perform precipitation control of TiC to finely precipitate in ferrite
at the time of austenite/ferrite transformation during cooling in the subsequent
second-stage cooling step and to contribute to strength. On the other hand,
the lower limit of the cooling stop temperature of the first-stage cooling step
20 is not limited in particular. However, a cooling stop temperature of the
subsequent second-stage cooling step to be performed for exhibiting
precipitation strengthening of ferrite is higher than 600°C as a condition of
exhibiting precipitation strengthening of ferrite. Thus, if the cooling stop
temperature of the first-stage cooling step is 600°C or lower, precipitation
25 strengthening cannot be obtained. Further, when it becomes an Arl point or
lower, ferrite cannot be obtained, to thus make it impossible to obtain an
40
intended microstructure.
[0120] In the second-stage cooling step to be performed next, an average
cooling rate is 10°C/sec or less, and in the present invention, air cooling
(standing-to-cool) is kept in mind. During cooling in this temperature zone,
5 transformation to ferrite from austenite is promoted, and simultaneously with
the transformation, fine carbide of Ti precipitates in ferrite, and an intended
strength of the steel sheet is obtained. When this cooling rate is greater than
10°C/sec, a moving speed of an interface between these two phases in the
transformation to ferrite from austenite becomes too fast, so that the
10 precipitation of carbide of Ti at the interface between the phases cannot keep
up with it and sufficient precipitation strengthening cannot be obtained.
[0121] Further, when it is greater than 10°C/sec, the transformation to
ferrite from austenite is delayed and an intended microstructure cannot be
obtained. On the other hand, the lower limit of the cooling rate in the
15 second-stage cooling step does not have to be limited in particular in the
present invention. However, unless heat input is performed externally by a
heating device, or the like, the cooling rate in the air cooling is 3°C/sec or so
even though the sheet thickness is half an inch or so, which is an upper limit
sheet thickness assumed in the present invention.
20 [0122] Further, a cooling time period in the second-stage cooling step is 1
second or longer and shorter than 100 seconds. This step is an extremely
important step not only for promoting two-phase separation of ferrite and
austenite to obtain an intended second phase fraction but also for promoting
precipitation strengthening by fine carbide of Ti in ferrite obtained after the
25 transformation being completed. When this time period is shorter than 1
second, the ferrite transformation does not progress and an intended
41
microstructure cannot be obtained, and furthermore the precipitation of
carbide of Ti in ferrite obtained after the transformation does not progress, so
that intended strength and burring workability of the steel sheet cannot be
obtained. When it is shorter than 3 seconds, the ferrite transformation and
5 the precipitation of carbide do not progress sufficiently, so that it is desirably
3 seconds or longer because there is a risk that low-temperature
transformation products and strength of ferrite are no longer able to be
obtained sufficiently.
[0123] On the other hand, even when it is 100 seconds or longer, the
10 above-described effect is saturated and further productivity decreases
significantly. When it is 15 seconds or longer, the average crystal diameter
of the low-temperature transformation product of the dual phase steel sheet
becomes coarse, and further there is a concern that pearlite is mixed into the
microstructure, so that it is desirably shorter than 15 seconds.
15 [0124] The cooling stop temperature in the second-stage cooling step is
higher than 600°C. When this temperature is 600°C or lower, the
precipitation of carbide of Ti in ferrite obtained after transformation does not
progress, so that the strength decreases.
[0125] On the other hand, the upper limit of the cooling stop temperature
20 in the second-stage cooling step is not defined in particular, but when it is
higher than 700°C, two-phase separation of ferrite and austenite is not
sufficient and an intended fraction of the low-temperature transformation
product cannot be obtained, and furthermore the precipitation of carbide of Ti
in ferrite is over-aged and the strength decreases.
25 [0126] In a third-stage cooling step to be subsequently performed,
cooling is performed at a cooling rate of 15°C/sec or more. When this
42
cooling rate is less than 15°C/sec, pearlite is mixed into the microstructure,
and thereby an intended microstructure might not be obtained. Incidentally,
a completing temperature of the third-stage cooling step is a coiling
temperature. Even though the upper limit of the cooling rate in the
5 third-stage cooling step is not limited in particular, the effect of the present
invention can be obtained, but when a sheet ward caused by thermal strain is
considered, it is desirably set to 300°C/sec or less.
[0127] In the third-stage cooling step, the steel sheet is cooled down to a
temperature zone of 350°C or lower to be coiled. When this temperature is
10 higher than 350°C, intended low-temperature transformation products cannot
be obtained. Concretely, coarse carbide is formed between laths of bainite
constituting the low-temperature transformation product to be a starting point
of occurrence of a crack at the time of burring, and the burring workability
deteriorates.
15 [0128] On the other hand, the lower limit value of the coiling temperature
does not have to be limited in particular, but when a coil is in a state of being
exposed to water for a long time, appearance failure caused by rust is
concerned, so that it is desirably 50°C or higher. Further, when this
temperature is 100°C or lower, most of the low-temperature transformation
20 product turns into fresh martensite and uniform elongation improves to be
advantageous to forming with, a dominant n value such as bulging.
[0129] In order to more efficiently exhibit precipitation strengthening by
carbide of Ti in the cooling step after finish rolling, it is necessary to control a
cooling pattern up to coiling itself. Concretely, a total cumulative diffusion
25 length Lt0(ai of Ti in ferrite expressed by Expression (2) below is controlled in
the range of not less than 0.15 nor more than 0.5.
43
[0130] That is, when the total cumulative diffusion length Ltotai of Ti in
ferrite is expressed by Expression (3) below by adding up a diffusion length L
of Ti in ferrite expressed by Expression (2) below for a very short time period
At/sec from a cooling completing temperature to coiling, 0.15 ^ L,otai ^
5 0.5 is satisfied.
L = VD(T + 273)t - Expression (2)
Ltotai = ZV(D(T + 273)At) - Expression (3)
[0131] Here, D(T + 273) is a volume diffusion coefficient at T°C and t is
a diffusion time period, and D(T) is expressed by Expression (4) below using
10 a diffusion coefficient DO of Ti, an activation energy Q, and a gas constant R.
D(T) = DO x Exp(-Q/R(T + 273)) - Expression (4)
[0132] When this L(ota] value is less than 0.15 jam, the precipitation of
carbide of Ti does not progress during cooling to result in underaging,
resulting in that precipitation strengthening ability cannot be obtained
15 efficiently. On the other hand, when it is greater than 0.5 jim, the
precipitation of carbide of Ti progresses too much during cooling to result in
overaging, resulting in that precipitation strengthening ability cannot be
obtained efficiently after all.
[0133] Incidentally, for the puipose of achieving improvement in ductility
20 by shape correction of the steel sheet and introduction of mobile dislocation,
skinpass rolling at a reduction ratio of not less than 0.1% nor more than 2% is
desirably performed after all the steps are completed. Further, for the
purpose of removing scales attached to the surface of an obtained hot-rolled
steel sheet, pickling may also be performed on the obtained hot-rolled steel
25 sheet according to need after all the steps are completed. Further, after the
pickling, on the obtained hot-rolled steel sheet, skinpass at a reduction ratio of
44
10% or less may also be performed inline or offline, or cold rolling at a
reduction ratio of down to 40% or so may also be performed.
[0134] Further, before or after, or before and after the skinpass rolling,
scales on the surface are removed. The step of removing scales is not
5 defined in particular. For example, general pickling using hydrochloric acid
or sulfuric acid, or a device according to a line such as surface grinding by a
sander or the like or surface scarfing using plasma, a gas burner, or the like
can be applied.
[0135] Further, after casting, after hot rolling, or after cooling, a hot
10 treatment may be performed on a hot-rolled steel sheet with the present
invention applied thereto on a hot-dip plating line, and further on the
hot-roiled steel sheet, a surface treatment may also be performed additionally.
Plating is performed on the hot-dip plating line, and thereby corrosion
resistance of the hot-rolled steel sheet improves.
15 [0136] Incidentally, when galvanizing is performed on the hot-rolled steel
sheet obtained after pickling, the obtained steel sheet may also be immersed
in a galvanizing bath to be subjected to an alloying treatment according to
need. By performing the alloying treatment, the hot-rolled steel sheet
improves in welding resistance against various weldings such as spot welding
20 in addition to the improvement in corrosion resistance.
Example
[0137] Steels A to Z and a to d having chemical components shown in
Table 1 were melted in a converter refining and secondaiy refining step, steel
billets (slabs) manufactured by continuous casting were each reheated and
25 reduced to a sheet thickness of 2.3 to 3.4 mm by finish rolling subsequently to
rough rolling, and were each cooled on a run-out-table to then be coiled, and
45
hot-rolled steel sheets were prepared. More specifically, in accordance with
manufacturing conditions shown in Tables 2 and 3, hot-rolled steel sheets
were prepared. Incidentally, chemical compositions in Table 1 all mean
mass%.
5 [0138] In Table 1, Ti* represents [Ti] - 48/14[N] - 48/32[S], in Tables 1
and 2, Ex.C represents [C] - 12/48 x ([Ti] + 48/93[Nb] - 48/14[N] - 48/32[S]),
and in Table 1, Mn/S represents [Mn]/[S]. Further, the balance of the
component in Table 1 is Fe and impurities, each underline in Tables 1 and 2
indicates that a numerical value is outside the range of the present invention.
10 Steels K and R each do not contain Si intentionally. In Table 1, "-" indicates
that no intentional containing is performed.
[0139] In Table 2, "STEEL" indicates a steel having the components
corresponding to each symbol shown in Table 1. "SOLUTION
TEMPERATURE" indicates the minimum slab reheating temperature (=
15 SRTmin) calculated by Expression (1). "Ar3 TRANSFORMATION POINT
TEMPERATURE" indicates a temperature calculated by Expression (5), (6),
or (7). "Ex.C" indicates a value calculated by [C] - 12/48 x ([Ti] +
48/93[Nb] - 48/14[N] - 48/32[S]).
[0140] In the manufacturing conditions in Tables 2 and 3, in the heating
20 step, "HEATING TEMPERATURE" indicates a maximum ultimate
temperature in slab reheating and "HOLDING TIME PERIOD" indicates a
holding time period at a predetermined heating temperature. In rough
rolling, "TOTAL PASS NUMBER" indicates a total value of the number of
rolling passes in rough rolling, "TOTAL REDUCTION RATIO" indicates a
25 reduction ratio in rough rolling from start to completion of rough rolling,
"NUMBER OF PASSES AT 1050 TO 1150°C AND AT 20% OR MORE"
46
indicates the number of passes of which rolling at a rolling ratio of 20% or
more was performed in a temperature zone of 1050 to 1150°C, "TIME
PERIOD UNTIL START OF FINISH ROLLING" indicates a time period
until start of finish rolling after completion of rough rolling, and "AVERAGE
5 AUSTENITE GRAIN DIAMETER IMMEDIATELY BEFORE FINISH
ROLLING" indicates an average grain diameter of austenite grains
immediately before a rough bar is bitten at the first stand of finish rolling.
Recognition of this austenite grain diameter can be obtained in a manner that
a crop piece obtained by cutting a rough bar before being subjected to finish
10 rolling by a crop shear or the like is quenched as much as possible to be
cooled down to room temperature or so, and a cross section parallel to a
rolling direction is etched to make austenite grain boundaries appear to
measure austenite grain diameters by an optical microscope. On this
occasion, 20 visual fields or more at the 1/4 position of a sheet thickness are
15 measured at 50 or more magnifications by an image analysis, a point counting
method, or the like.
[0141] In finish rolling, "ROLLING START TEMPERATURE" indicates
a temperature immediately before a rough bar is bitten at the first stand of
finish rolling, "TOTAL REDUCTION RATIO" indicates a reduction ratio
20 during finish rolling from start to completion of finish rolling, "AVERAGE
REDUCTION RATIO FOR 3 PASSES AT REAR STAND" indicates an
average value of reduction ratios from the final pass including the final pass
to the third pass in finish rolling in which continuous rolling with plural
passes is normally performed, "FINISH ROLLING OUTLET SIDE SPEED"
25 indicates an outlet side sheet passing speed at the rolling stand after a finish
rolling final reduction pass is completed, and "FINISHING
47
TEMPERATURE" indicates a temperature immediately after a rolling stand
outlet side of a finish roiling final pass. Incidentally, the reduction ratio may
be an actual performance value calculated from a sheet thickness, or may also
be a setup value of a rolling stand. Further, the temperature is desirably
5 measured at the step position with a radiation thermometer or a contact
thermometer, but may also be an estimated value obtained by a temperature
model or the like.
[0142] The cooling step performed on a run-out-table is divided into first
to third-stage cooling steps in terms of precipitation control and structure
10 control. First, in "FIRST-STAGE COOLING STEP," "TIME PERIOD
UNTIL START OF COOLING" indicates a time period until start of cooling
on a run-out-table after passing through a rolling stand of a finish rolling final
pass, "COOLING RATE" indicates an average cooling rate by water cooling,
and "COOLING STOP TEMPERATURE" indicates a temperature at which
15 water cooling in the first-stage cooling step is stopped. In
"SECOND-STAGE COOLING STEP," "COOLING RATE" indicates an
average cooling rate by air cooling without pouring water mainly,
"HOLDING TIME PERIOD" indicates a holding time period of air cooling
without pouring water, and "COOLING STOP TEMPERATURE" indicates a
20 temperature at which holding of air cooling without pouring water is
completed. In "THIRD-STAGE COOLING STEP," "COOLING RATE"
indicates an average cooling rate until restart of water cooling and coiling
after air cooling and holding, and "COILING TEMPERATURE" indicates a
temperature immediately before a steel sheet is coiled into a coil shape by a
25 coiler after stopping water cooling. Incidentally, "TOTAL CUMULATIVE
DIFFUSION LENGTH" indicates the total cumulative diffusion length Ltotal
48
of Ti in ferrite and is obtained by Expression (3) by adding up the diffusion
length L of Ti in ferrite expressed by Expression (2) for the very short time
period At/sec from a cooling completing temperature to coiling.
[0143] Microstructures of steel sheets obtained by manufacturing
5 methods described in Tables 2 and 3 are shown in Table 4, and mechanical
property, surface property, and corrosion resistance are shown in Table 5.
[0144] First, a sample was taken from the 1/4W position or 3/4W position
of a sheet width of each of the obtained steel sheets, and by using an optical
microscope, each microstructure at the 1/4 thickness of a sheet thickness was
10 observed. As adjustment of the samples, a sheet thickness cross section in
the rolling direction was polished as an observation surface to be subjected to
etching with a nital reagent and a LePera reagent. From each optical
micrograph at 500 magnifications of the sheet thickness cross sections etched
with a nital reagent and a LePera reagent, "MICROSTRUCTURE" was
15 classified.
[0145] Further, from each of optical micrographs at 500 magnifications
of the sheet thickness cross sections etched with a LePera reagent, "SECOND
PHASE CHARACTERISTIC" being a distribution state of the
low-temperature transformation product being a second phase was recognized
20 by an image analysis. Here, the dispersion state of the low-temperature
transformation product is classified into one in which the low-temperature
transformation products are dispersed in an island shape at a corner, an edge,
and a grain boundary surface of a ferrite grain as "ISLAND SHAPE," one in
which they are island shaped but are distributed continuously parallel to the
25 rolling direction as "ALIGNED STATE," and one in which they are dispersed
to surround a grain boundary surface of a ferrite grain mainly as "FILM
49
SHAPE."
[0146] Further, by the image analysis, "SECOND PHASE FRACTION"
being the area fraction of the low-temperature transformation product being a
second phase and "SECOND PHASE AVERAGE GRAIN DIAMETER"
5 being the average grain diameter of the low-temperature transformation
product were obtained. "Ex.C (%)/fsd (%)" is a value of "Ex.C (%)" in
Table 2 divided by "SECOND PHASE FRACTION." Incidentally, the
average crystal diameter of the low-temperature transformation product is one
in which circle-equivalent diameters are number-averaged. Further, plural
10 low-temperature transformation products were selected arbitrarily, respective
distances of closest approach of them were obtained, and an average value of
20 points was set to "AVERAGE VALUE OF DISTANCE OF CLOSEST
APPROACH BETWEEN SECOND PHASES."
[0147] Nanohardness Hn was measured by using
15 TriboScope/TriboIndenter manufactured by Hysitron. As the measuring
condition, hardness of the low-temperature transformation product was
measured at 20 points or more with I mN of load, and an arithmetic average
of them and a standard deviation were calculated.
[0148] Measurement of "FERRITE TiC DENSITY" being a TiC
20 precipitate density was performed by a three-dimensional atom probe
measurement method. First, an acicular sample is prepared from a sample to
be measured by cutting and electropolishing, and by using focused ion beam
milling together with electropolishing according to need. In the
three-dimensional atom probe measurement, integrated data can be
25 reconstructed to obtain an actual distribution image of atoms in a real space.
A number density of TiC precipitates is obtained from the volume of a
50
three-dimensional distribution image of TiC precipitates and the number of
TiC precipitates. Incidentally, the measurement was performed in a manner
that ferrite grains are specified and five or more of ferrite grains for each
sample are used. Further, as for the size of the above-described TiC
5 precipitates, a diameter calculated from the number of atoms constituting
observed TiC precipitates and a lattice constant of TiC assuming that the
precipitates are spherical is set as the size. Arbitrarily, diameters of 30 or
more of TiC precipitates were measured. An average value of them was 2 to
30 nm or so.
10 [0149] Of the mechanical property, tensile strength properties (YP, TS,
and El) were evaluated based on JIS Z 2241-1998 by using a No. 5 test piece
of JIS Z 2201-1998 taken from the 1/4W position or the 3/4W position of the
sheet width in a direction vertical to the rolling direction. As an index of the
burring workability, a hole expanding test was employed. With regard to the
15 hole expanding test, a test piece was taken from the same position as that
where a tensile test piece was taken, and the burring workability was
evaluated based on a test method described in Japan Iron and Steel Federation
specification JFS T 1001-1996.
[0150] Next, in order to examine the notch fatigue strength, a fatigue test
20 piece having a shape shown in FIG. 1 was taken from the same position as
that where the tensile test piece was taken so that the side in the rolling
direction could be a long side and was subjected to a fatigue test. Here, the
fatigue test piece described in FIG. 1 is a notched test piece prepared for
obtaining the notch fatigue strength. Side surface corner portions (portions
25 each surrounded by a dotted line in FIG. 1) of this notched test piece are each
chamfered with 1R to be polished in the longitudinal direction with #600.
51
[0151] In order to approach the fatigue property evaluation in actual use
of an automobile part, the notch was made by punching with a cylinder punch
in the same manner as that of the hole expanding test piece. Incidentally, a
punching clearance was set to 12.5%. However, on the fatigue test piece,
5 grinding of fine finishing was performed down to the depth of 0.05 mm or so
from the uppermost surface layer. A Schenck type fatigue testing machine
was used for the fatigue test, and a test method was based on JIS Z 2273-1978
and JIS Z 2275-1978. "awk/TS" being the definition of the notch fatigue
property in Table 3 is a value of a 2 million cycle fatigue strength obtained by
10 this test divided by a tensile strength.
[0152] The surface property was evaluated by "SURFACE DEFECT"
and "ROUGHNESS" before pickling. When this evaluation is equal to or
less than the reference, there is sometimes a case that the surface quality is
evaluated according to a pattern and unevenness of the surface caused by a
15 scale defect by inferiors and customers even after pickling. Here,
"SURFACE DEFECT" indicates a result obtained by visually recognizing the
presence/absence of scale defects such as Si scales, scales, and spindles, and
the case of scale defects being present is shown as "x" and the case of no
scale defects is shown as "O-" Incidentally, one in which these defects are
20 partial or the reference or less is regarded as "SLIGHT" to be shown as "A."
"ROUGHNESS" is evaluated by Rz and indicates a value obtained by a
measurement method described in JIS B 0601-2001. Incidentally, as long as
Rz is 20 um or less, the surface quality is a level with no problems.
[0153] The corrosion resistance was evaluated by "CONVERSION
25 TREATMENT PROPERTY" AND "POST-COATING CORROSION
RESISTANCE." First, the manufactured steel sheet was pickled, and then
52
was subjected to a conversion treatment in which a zinc phosphate coating
film of 2.5 g/m is attached. At this stage, measurements of
presence/absence of lack of hiding and a P ratio were performed as
"CONVERSION TREATMENT PROPERTY."
5 [0154] The phosphoric acid conversion treatment is a treatment using a
chemical solution having phosphoric acid and Zn ions as its main component,
and is a chemical reaction to generate a crystal called phosphophyllite:
FeZn2(P04)3-4H20 between Fe ions to liquate from the steel sheet. The
technical points of the phosphoric acid conversion treatment are to (1) make
10 Fe ions liquate to promote the reaction and to (2) densely form
phosphophyllite crystals on the surface of the steel sheet. Particularly, with
regard to (1), when oxides ascribable to formation of Si scales remain on the
surface of the steel sheet, liquation of Fe is prevented and a portion to which a
conversion coating film does not attach, which is called lack of hiding,
15 appears, due to no liquation of Fe, an abnormal conversion treatment coating
film that is not formed normally on the surface of an iron, called hopeite:
Zn3(P04)3-4H20, is formed, and thereby performance after coating
sometimes deteriorates. Thus, it becomes important to make the surface
normal so that by liquating Fe on the surface of the steel sheet by phosphoric
20 acid, Fe ions can be supplied sufficiently.
[0155] This lack of hiding can be recognized by observation by a
scanning electron microscope, 20 visual fields or so are observed at 1000
magnifications, and the case where the conversion coating film is uniformly
attached to the entire surface and no lack of hiding can be recognized is
25 regarded as no lack of hiding to be shown as "O." Further, the case where
the visual field with recognition of lack of hiding is 5% or less is regarded as
53
slight to be shown as "A." Further, the case where it is greater than 5% is
regarded as presence of lack of hiding to be evaluated as "x."
[0156] On the other hand, the P ratio can be measured by using an X-ray
diffraction device, a ratio of an X-ray diffraction intensity P of the
5 phosphophylhte (100) plane and an X-ray diffraction intensity H of the
hopeite (020) plane is taken, and the P ratio is evaluated by P ratio = P/(P + H).
That is, the P ratio represents the ratio of hopeite and phosphophylhte in the
coating film obtained by performing the conversion treatment, and it means
that as the P ratio is higher, phosphophylhte is more contained and
10 phosphophylhte ciystals are densely formed on the surface of the steel sheet.
Generally, P ratio ^ 0.80 is required in order to satisfy anticorrosion
performance and coating performance, and under severe corrosive
environment such as in a thawing salt scattering region, P ratio ^ 0.85 is
required.
15 [0157] Next, with regard to the corrosion resistance, electrodeposition
coating to have a thickness of 25 (am was performed after the conversion
treatment and a coating and baking treatment at 170°C x for 20 minutes was
performed, and then an incision having a length of 130 mm was made in an
electrodeposition coating film to reach the base iron with a knife having a
20 sharp end, and under a salt spray condition described in JIS Z 2371, 5% salt
spraying at a temperature of 35°C was performed for 700 hours continuously
and then a tape (Nichiban Co., Ltd. 405A-24 JIS Z 1552) having a width of
24 mm and having a length of 130 mm was applied on the incision portion
parallel to the incision portion, and the maximum coating film peeled width
25 obtained after the tape was peeled off was measured. This maximum
coating film peeled width of greater than 4 mm was defined that the corrosion
54
resistance is inferior.
[0158] Next, results will be explained. Incidentally, with regard to Steel
numbers 32, 36, and 46, the sheet was passed through an alloying hot-dip
galvanizing line after the pickling, and at a Zn bath temperature of 430 to
5 460°C, plating bath immersion was performed, and on Steel 32 and 46 out of
them, an alloying treatment was further performed at an alloying temperature
of500to600°C.
[0159] Steel numbers 1, 4, 9, 10, 11, 20, 23, 24, 25, 26, 27, 28, 29, 30, 31,
32, 33, 34, 35, 36, 37, 38, and 39 are in accordance with the present invention.
10 These steel sheets are steel sheets of grades being 540 MPa grade and
higher that contain predetermined amounts of steel components and in which
at the position of 1/4 thickness of the sheet thickness, a microstructure is a
dual phase with its main phase composed of polygonal ferrite
precipitation-strengthened by carbide of Ti and its second phase composed of
15 1 to 10% in area fraction (fsd (%)) of low-temperature transformation
products dispersed in an island shape, 0.001 ^ Ex.C (%)/fsd (%) ^ 0.01
(Ex.C (%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]})
is satisfied, an average crystal diameter of the low-temperature transformation
product is 3 to 15 um, and an average value of a distance of closest approach
20 between the low-temperature transformation products is 10 to 20 um, and
high-strength steel sheets having a hole expansion value X ^ 70%, having a
notch fatigue property of aWK/TS ^ 0.35, and having slight surface defects
or no surface defects can be obtained.
Steel numbers 32 and 39 contain Steel K and R containing no Si
25 intentionally respectively, and the content of Si of them is 0 or an impurity
level. However, Steel numbers 32 and 39 also satisfy the mechanical
55
property of the present invention.
[0160] The steels other the above are outside the range of the present
invention due to the following reasons.
That is, with regard to Steel number 2, the heating temperature is
5 outside the range of the manufacturing method of the present invention steel,
so that the predetermined microstructure cannot be obtained and the tensile
strength is low.
With regard to Steel number 3, the total reduction ratio of the rough
rolling is outside the range of the manufacturing method of the present
10 invention steel, so that the predetermined microstructure cannot be obtained
and the hole expansion value is low.
With regard to Steel number 5, the number of passes at 1050 to
1150°C and at 20% or more is outside the range of the manufacturing method
of the present invention steel, so that the predetermined microstructure cannot
15 be obtained and the hole expansion value is low.
[0161] With regard to Steel number 6, the time period until start of the
finish rolling is outside the range of the manufacturing method of the present
invention steel, so that the predetermined microstructure cannot be obtained
and the tensile strength and the hole expansion value are low.
20 With regard to Steel number 7, the finish rolling start temperature is
outside the range of the manufacturing method of the present invention steel,
so that the predetermined microstructure cannot be obtained and the tensile
strength is low.
With regard to Steel number 8, the total reduction ratio of the finish
25 rolling is outside the range of the manufacturing method of the present
invention steel, so that the predetermined microstructure cannot be obtained
56
and the hole expansion value is low.
With regard to Steel number 12, the. finish rolling finishing
temperature is outside the range of the manufacturing method of the present
invention steel, so that the predetermined microstructure cannot be obtained
5 and the hole expansion value is low.
With regard to Steel number 13, the finish rolling finishing
temperature is outside the range of the manufacturing method of the present
invention steel, so that the predetermined microstructure cannot be obtained
and the hole expansion value is low.
10 [0162] With regard to Steel number 14, the time period until the cooling
is outside the range of the manufacturing method of the present invention
steel, so that the predetermined microstructure cannot be obtained and the
tensile strength and the hole expansion value are low.
With regard to Steel number 15, the cooling rate of the cooling (a) is
15 outside the range of the manufacturing method of the present invention steel,
so that the predetermined microstructure cannot be obtained and the hole
expansion value and the notch fatigue property are low.
With regard to Steel number 16, the cooling stop temperature of the
cooling (a) is outside the range of the manufacturing method of the present
20 invention steel, so that the predetermined microstiiicture cannot be obtained
and the tensile strength and the notch fatigue property are low.
With regard to Steel number 17, the cooling stop temperature of the
cooling (a) is outside the range of the manufacturing method of the present
invention steel, so that the predetermined microstructure cannot be obtained
25 and the tensile strength and the notch fatigue property are low.
[0163] With regard to Steel number 18, the cooling rate of the cooling (b)
57
is outside the range of the manufacturing method of the present invention
steel, so that the predetermined microstructure cannot be obtained and the
tensile strength and the hole expansion value are low.
With regard to Steel number 19, the holding time period of the cooling
5 (b) is outside the range of the manufacturing method of the present invention
steel, so that the predetermined microstructure cannot be obtained and the
tensile strength and the notch fatigue property are low.
With regard to Steel number 21, the cooling rate of the cooling (c) is
outside the range of the manufacturing method of the present invention steel,
10 so that the predetermined microstructure cannot be obtained and the hole
expansion value and the notch fatigue property are low.
With regard to Steel number 22, the coiling temperature is outside the
range of the manufacturing method of the present invention steel, so that the
predetermined microstructure cannot be obtained and the hole expansion
15 value is low.
[0164] With regard to Steel number 40, the content of C is outside the
range of the present invention steel, so that the predetermined microstructure
cannot be obtained and the hole expansion value is low.
With regard to Steel number 41, the content of C is outside the range
20 of the present invention steel, so that the predetermined microstructure cannot
be obtained and the tensile strength is low.
With regard to Steel number 42, the content of Si is outside the range
of the present invention steel, so that the surface property is poor.
With regard to Steel number 43, the content of Mn is outside the range
25 of the present invention steel, so that slab cracking occurs to make the rolling
impossible.
58
With regard to Steel number 44, the content of Mn is outside the range
of the present invention steel, so that the predetermined microstructure cannot
be obtained and the tensile strength is low.
[0165] With regard to Steel number 45, the content of P is outside the
5 range of the present invention steel, so that the elongation and the notch
fatigue property are low due to embrittlement.
With regard to Steel number 46, the content of S is outside the range
of the present invention steel, so that MnS becomes a starting point of a crack
and the hole expansion value is low.
10 With regard to Steel number 47, the content of N is outside the range
of the present invention steel, so that coarse TiN becomes a starting point of a
crack and the hole expansion value is low.
[0166] With regard to Steel number 48, the content of Ti is outside the
range of the present invention steel, so that the predetermined microstructure
15 cannot be obtained and the notch fatigue property is low.
With regard to Steel number 49, the content of Ti is outside the range
of the present invention steel, so that the predetermined microstructure cannot
be obtained and the tensile strength is low.
With regard to Steel number 50, the value of Ti* is outside the range
20 of the present invention steel, so that the predetermined microstructure cannot
be obtained and the hole expansion value and the notch fatigue property are
low.
With regard to Steel number 51, the content of Al is outside the range
of the present invention steel, so that the predetermined microstructure cannot
25 be obtained and the hole expansion value is low.
64
[Industrial Applicability]
[0172] The dual phase steel sheet of the present invention can be used for
various uses such as shipbuilding, construction, bridges, offshore structures,
pressure vessels, linepipes, and machine parts, in addition to automobile
5 members that are required to have workability, hole expandability, and
bendability as well as having high strength such as inner sheet members,
structure members, and underbody members.
[Name of Document] Claim
[Claim 1] A dual phase steel sheet comprising:
in mass%,
C: 0.01 to 0.1%;
Mn: 0.2 to 3%;
Al: 0.04 to 1.5%;
Ti: 0.015 to 0.2%;
Si: 0 to 0.5%;
Nb:0to0.06%;
Cu:0tol.2%;
Ni: 0 to 0.6%;
Mo:0to 1%;
V:0to0.2%;
Cr: 0 to 2%;
W:0to0.5%;
Mg:0to0.01%;
Ca:0to0.01%;
REM: 0 to 0.1%;
B:0 to 0.002%;
P: 0.01% or less;
S: 0.005% or less;
N: 0.01% or less,
in which [Ti] - 48/14 x [N] - 48/32 x [S] ^ 0% is satisfied and when Ex.C
(%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]} is set,
0.001 ^ Ex.C (%)/fsd (%) £ 0.01 is satisfied, and
a balance being composed of Fe and impurities, wherein
at the position of 1/4 thickness of a sheet thickness, a microstructure is
a dual phase with its main phase composed of polygonal ferrite
precipitation-strengthened by carbide of Ti and its second phase composed of
1 to 10% in area fraction (fsd (%)) of low-temperature transformation
products dispersed plurally, and
an average crystal diameter of the low-temperature transformation
product is 3 to 15 um and an average value of a distance of closest approach
between the low-temperature transformation products is 10 to 20 um.
[Claim 2] The dual phase steel sheet according to claim 1, comprising:
in mass%,
Si: 0.02% to 0.5%.
[Claim 3] The dual phase steel sheet according to claim 1 or 2,
comprising:
one or two or more of
in mass%,
Nb: 0.005 to 0.06%;
Cu: 0.02 to 1.2%;
Ni: 0.01 to 0.6%;
Mo: 0.01 to 1%;
V:0.01 to 0.2%;
Cr:0.01 to 2%; and
W: 0.01 to 0.5%.
[Claim 4] The dual phase steel sheet according to any one of claims 1 to
3, comprising:
one or two or more of
in mass%,
Mg: 0.0005 to 0.01%;
Ca: 0.0005 to 0.01%; and
REM: 0.0005 to 0.1%.
[Claim 5] The dual phase steel sheet according to any one of claims 1 to
4, comprising:
in mass%,
B: 0.0002 to 0.002%.
[Claim 6] The dual phase steel sheet according to any one of claims 1 to
5, wherein
galvanizing is performed on its surface.
[Claim 7] A manufacturing method of a dual phase steel sheet
comprising:
on a slab containing:
in mass%,
C: 0.01 to 0.1%;
Mn: 0.2 to 3%;
Al: 0.04 to 1.5%;
Ti: 0.015 to 0.2% or less;
Si: 0 to 0.5%;
Nb:0to0.06%;
Cu:0tol.2%;
Ni:0to0.6%;
Mo:0tol%;
V: 0 to 0.2%;
Cr: 0 to 2%;
W:0to0.5%;
68
Mg:0to0.01%;
Ca:0to0.01%;
REM: 0 to 0.1%;
B:0 to 0.002%;
P: 0.01% or less;
S: 0.005% or less;
N: 0.01% or less,
in which [Ti] - 48/14 x [N] - 48/32 x [S] ^ 0% is satisfied and when Ex.C
(%) = [C] - 12/48 x {[Ti] + 48/93 x [Nb] - 48/14 x [N] - 48/32 x [S]} is set,
0.001 ^ Ex.C (%)/fsd (%) ^ 0.01 is satisfied, and
a balance being composed of Fe and impurities, performing heating to
a temperature SRTmin (°C) or higher, which is defined by Expression (1)
below, and then in hot roiling, performing rough rolling at a reduction ratio of
20%o or more in a temperature zone of not lower than 1050°C nor higher than
1150°C for at least one pass, and then starting finish rolling within 150
seconds in a temperature zone of 1000°C or higher and lower than 1080°C,
and completing finish rolling with the total reduction ratio for plural passes of
not less than 75% nor more than 95% in a temperature zone of not lower than
an Ar3 transformation point temperature + 50°C nor higher than 1000°C; and
within 3 seconds, performing cooling down to lower than the Ar3
transformation point temperature at an average cooling rate of 15°C/sec or
more, and next performing cooling down to a temperature zone of higher than
600°C at an average cooling rate of 10°C/sec or less for a time period of 1
second or longer and shorter than 300 seconds, and next performing cooling
down to a temperature zone of 350DC or lower at a cooling rate of 15°C/sec or
more, and performing coiling.
69
SRTmin = 10780/(5.13 - log([Ti] x [C])} - 273 - Expression (1)
[Claim 8] The manufacturing method of the dual phase steel sheet
according to claim 7, further comprising:
in the hot rolling, performing rough rolling at a reduction ratio of 20%
or more in a temperature zone of not lower than 1050°C nor higher than
1150°C for plural passes, wherein
the total reduction ratio of the rough rolling is not less than 60% nor
more than 90%.
[Claim 9] The manufacturing method of the dual phase steel sheet
according to claim 7 or 8, further comprising:
performing cooling down to a temperature zone of 100°C or lower and
performing coiling.
[Claim 10] The manufacturing method of the dual phase steel sheet
according to any one of claims 7 to 9, wherein
in the performing the cooling down to the temperature zone of higher
than 600°C at an average cooling rate of 10°C/sec or less for a time period of
1 second or longer and shorter than 100 seconds, when a total cumulative
diffusion length Ltotai of Ti in ferrite is expressed by Expression (3) below by
adding up a diffusion length L of Ti in ferrite expressed by Expression (2)
below for a very short time At/sec from a cooling completing temperature to
coiling, 0.15 ^ Llotai ^ 0.5 is satisfied.
L = VD(T + 273)t - Expression (2)
Ltotai = SV(D(T + 273)At) - Expression (3)
Here, D(T + 273) is a volume diffusion coefficient at T°C. t is a
diffusion time period.
D(T) is expressed by Expression (4) below using a diffusion coefficient DO of
Ti, an activation energy Q, and a gas constant R.
D(T) = DO x Exp(-Q/R-(T + 273)) - Expression (4)
[Claim 11] The manufacturing method of the dual phase steel sheet
according to any one of claims 7 to 10, wherein
in the performing the cooling down to the temperature zone of higher
than 600°C at an average cooling rate of 10°C/sec or less for a time period of
1 second or longer and shorter than' 100 seconds, a steel sheet is immersed in
a galvanizing bath to galvanize its surface.
[Claim 12] The manufacturing method of the dual phase steel sheet
according to claim 11, further comprising:
on a galvanized dual phase steel sheet, performing an alloying
treatment in a temperature range of 450 to 600°C.