STEEL MATERIAL
[Technical Field]
[0001] The present invention relates to a steel material, and concretely
5 relates to a steel material suitable for a material of an impact absorbing
member in which an occurrence of crack when applying an impact load is
suppressed, and further, an effective flow stress is high. This application is
based upon and claims the benefit of priority of the prior Japanese Patent
Application No. 2012-182710, filed on August 21, 2012, the entire contents of
10 which are incorporated herein by reference.
[Background Art]
[0002] In recent years, from a point of view of global environmental
protection, a reduction in weight of a vehicle body of automobile has been
required as a part of reduction in C02 emissions from automobiles, and a
15 high-strengthening of a steel material for automobile has been aimed. This
is because, by improving the strength of steel material, it becomes possible to
reduce a thickness of the steel material for automobile. Meanwhile, a social
need with respect to an improvement of collision safety of automobile has
been further increased, and not only the high-strengthening of steel material
20 but also a. development of steel material excellent in impact resistance when a
collision occurs during traveling, has been desired.
[0003] Here, respective portions of a steel material for automobile at a
time of collision are deformed at a high strain rate of several tens (s"1) or more,
so that a high-strength steel material excellent in dynamic strength property is
25 required.
[0004] As such a high-strength steel material, a low-alloy TRIP steel
2
having a large static-dynamic difference (difference between static strength
and dynamic strength), and a high-strength multi-phase structure steel
material such as a multi -phase structure steel having a second phase mainly
formed of martensite, are known.
5 [0005] Regarding the low-alloy TRIP steel, for example, Patent
Document 1 discloses a strain-induced transformation type high-strength steel
sheet (TRIP steel sheet) for absorbing collision energy of automobile
excellent in dynamic deformation property.
[0006] Further, regarding the multi-phase structure steel sheet having the
10 second phase mainly formed of martensite, inventions as will be described
below are disclosed.
[0007] Patent Document 2 discloses a high-strength steel sheet having an
excellent balance of strength and ductility and having a static-dynamic
difference of 170 MPa or more, the high-strength steel sheet being formed of
15 fine ferrite grains, in which an average grain diameter ds of nanocrystal grains
each having a crystal grain diameter of 1.2 jum or less and an average crystal
grain diameter dL of microcrystal grains each having a crystal grain diameter
of greater than 1.2 pm satisfy a relation of dL / ds ^ 3.
[0008] Patent Document 3 discloses a steel sheet formed of a dual-phase
20 structure of martensite whose average grain diameter is 3 urn or less and
martensite whose average grain diameter is 5 u,m or less, and having a high
static-dynamic ratio.
[0009] Patent Document 4 discloses a cold-rolled steel sheet excellent in
impact absorption property containing 75% or more of ferrite phase in which
25 an average grain diameter is 3.5 u,m or less, and a balance composed of
tempered martensite.
3
[0010] Patent Document 5 discloses a cold-rolled steel sheet in which a
prestrain is applied to produce a dual-phase structure formed of ferrite and
martensite, and a static-dynamic difference at a strain rate of 5 x 102 to 5 x
103/s satisfies 60 MPa or more.
5 [0011] Further, Patent Document 6 discloses a high-strength hot-rolled
steel sheet excellent in impact resistance property formed only of hard phase
such as bainite of 85% or more and martensite.
[Prior Art Document]
[Patent Document]
10 [0012] Patent Document 1: Japanese Laid-open Patent Publication No.
HI 1-80879
Patent Document 2: Japanese Laid-open Patent Publication No.
2006-161077
Patent Document 3: Japanese Laid-open Patent Publication No.
15 2004-84074
Patent Document 4: Japanese Laid-open Patent Publication No.
2004-277858
Patent Document 5: Japanese Laid-open Patent Publication No.
2000-17385
20 Patent Document 6: Japanese Laid-open Patent Publication No.
HI 1-269606
[Disclosure of the Invention]
[Problems to Be Solved by the Invention]
[0013] However, the conventional steel materials being materials of
25 impact absorbing members have the following problems. Specifically, in
order to improve an impact absorption energy of an impact absorbing member
4
(which is also simply referred to as "member", hereinafter), it is essential to
increase a strength of a steel material being a material of the impact absorbing
member (which is also simply referred to as "steel material", hereinafter).
[0014] Incidentally, as disclosed in "Journal of the Japan Society for
5 Technology of Plasticity" vol. 46, No. 534, pages 641 to 645, that an average
load (Fave) determining an impact absorption energy is given in a manner that
Fave00 (cY • t2) / 4, in which aY indicates an effective flow stress, and t
indicates a sheet thickness, the impact absorption energy greatly depends on
the sheet thickness of steel material. Therefore, there is a limitation in
10 realizing both of a reduction in thickness and a high impact absorbency of the
impact absorbing member only by increasing the strength of the steel
material.
[0015] Here, the flow stress corresponds to a stress required for
successively causing a plastic deformation at a start or after the start of the
15 plastic deformation, and the effective flow stress means a plastic flow stress
which takes a sheet thickness and a shape of the steel material and a rate of
strain applied to a member when an impact is applied into consideration.
[0016] Meanwhile, for example, as disclosed in pamphlet of International
Publication No. WO 2005/010396, pamphlet of International Publication No.
20 WO 2005/010397, and pamphlet of International Publication No. WO
2005/010398, an impact absorption energy of an impact absorbing member
also greatly depends on a shape of the member.
[0017] Specifically, by optimizing the shape of the impact absorbing
member so as to increase a plastic deformation workload, there is a possibility
25 that the impact absorption energy of the impact absorbing member can be
dramatically increased to a level which cannot be achieved only by increasing
5
the strength of the steel material.
[0018] However, even when the shape of the impact absorbing member is
optimized to increase the plastic deformation workload, if the steel material
has no deformability capable of enduring the plastic deformation workload, a
5 crack occurs on the impact absorbing member in an early stage before an
expected plastic deformation is completed, resulting in that the plastic
deformation workload cannot be increased, and it is not possible to
dramatically increase the impact absorption energy. Further, the occurrence
of crack on the impact absorbing member in the early stage may lead to an
10 unexpected situation such that another member disposed by being adjacent to
the impact absorbing member is damaged.
[0019] In the conventional techniques, it has been aimed to increase the
dynamic strength of the steel material based on a technical idea that the
impact absorption energy of the impact absorbing member depends on the
15 dynamic strength of the steel material, but, there is a case where the
deformability is significantly lowered only by aiming the increase in the
dynamic strength of the steel material. Accordingly, even if the shape of the
impact absorbing member is optimized to increase the plastic deformation
workload, it was not always possible to dramatically increase the impact
20 absorption energy of the impact absorbing member.
[0020] Further, since the shape of the impact absorbing member has been
studied on the assumption that the steel material manufactured based on the
above-described technical idea is used, the optimization of the shape of the
impact absorbing member has been studied, from the first, based on the
25 deformability of the existing steel material as a premise, and thus the study
itself such that the deformability of the steel material is increased and the
6
shape of the impact absorbing member is optimized to increase the plastic
deformation workload, has not been done sufficiently so far.
[0021] The present invention has a task to provide a steel material
suitable for a material of an impact, absorbing member having a high effective
5 flow stress and thus having a high impact absorption energy and in which an
occurrence of crack when an impact load is applied is suppressed, and a
manufacturing method thereof.
[Means for Solving the Problems]
[0022] As described above, in order to increase the impact absorption
10 energy of the impact absorbing member, it is important to optimize not only
the steel material but also the shape of the impact absorbing member to
increase the plastic deformation workload.
[0023] Regarding the steel material, it is important to increase the
effective flow stress to increase the plastic deformation workload while
15 suppressing the occuiTence of crack when the impact load is applied, so that
the shape of the impact absorbing member capable of increasing the plastic
deformation workload can be optimized.
[0024] The present inventors conducted earnest studies regarding a
method of suppressing the occurrence of crack when the impact load is
20 applied and increasing the effective flow stress regarding the steel material to
increase the impact absorption energy of the impact absorbing member, and
obtained new findings as will be cited hereinbelow.
[Improvement of impact absorption energy]
(1) In order to increase the impact absorption energy of the steel
25 material, it is effective to increase the effective flow stress when a true strain
of 5% is given (which will be described as " 5% flow stress", hereinafter).
7
[0025] (2) In order to increase the 5% flow stress, it is effective to
increase a yield strength and a work hardening coefficient in a low-strain
region.
[0026] (3) In order to increase the yield strength, it is effective to produce
5 a steel structure containing bainite as a main phase.
[0027] (4) In order to increase the work hardening coefficient in the
low-strain region in the steel material containing bainite as the main phase, it
is effective to make fine precipitates exist at a high density.
[0028] [Suppression of occurrence of crack when impact load is applied]
10 (5) When a crack occurs on the impact absorbing member at the time
of applying the impact load, the impact absorption energy is lowered.
Further, there is also a case where another member adjacent to the impact
absorbing member is damaged.
[0029] (6) When the strength, particularly the yield strength of the steel
15 material is increased, a sensitivity with respect to a crack at the time of
applying the impact load (which is also referred to as "impact crack",
hereinafter) (the sensitivity is also referred to as "impact crack sensitivity",
hereinafter) becomes high.
[0030] (7) In order to suppress the occurrence of impact crack, it is
20 effective to increase a uniform ductility, a local ductility and a fracture
toughness.
[0031] (8) In the steel material containing bainite as the main phase, the
ductility can be increased by refining bainite being the main phase.
[0032] (9) It is set that the steel material containing bainite as the main
25 phase contains, as a second phase, one or two or more selected from a group
consisting of ferrite, martensite and austenite, and if the above elements are
8
refined, the local ductility can be further improved.
[0033] (10) In order to increase the fracture toughness in the steel
material containing bainite as the main phase, it is effective to produce a
structure in which ferrite is contained in the second phase. However, coarse
5 ferrite causes a decrease in the yield stress and a crush load, so that ferrite has
to be refined.
[0034] (11) In order to increase the uniform ductility in the steel material
containing bainite as the main phase, it is effective to produce a structure in
which austenite is contained in the second phase. However, coarse austenite
10 exerts an adverse effect on the fracture toughness when being transformed
into a martensite phase due to a strain induction, so that austenite has to be
refined.
[0035] (12) In order to increase the fracture toughness in the steel
material containing bainite as the main phase, it is effective to produce a
15 structure in which martensite is contained in the second phase. However,
coarse martensite exerts an adverse effect on the fracture toughness, so that
martensite has to be refined.
[0036] The present invention is made based on the above-described new
findings, and a gist thereof is as follows.
20 [0037] . [1]
A steel material contains: by mass%, C: greater than 0.05% to 0.18% ;
Mm 1% to 3% ; Si: greater than 0.5% to 1.8% ; Al: 0.01% to 0.5% ; N:
0.001% to 0.015% ; one or both of V and Ti: 0.01% to 0.3% in total ; Cr: 0%
to 0.25% ; Mo: 0% to 0.35% ; a balance: Fe and impurities; and 80%> or more
25 of bainite by area%, and 5% or more in total of one or two or more selected
from a group consisting of ferrite, martensite and austenite by area%, in
9
which an average block size of the above-described bainite is less than 2.0 ixm,
an average grain diameter of all of the above-described ferrite, martensite and
austenite is less than 1.0 urn, an average nanohardness of the above-described
bainite is 4.0 GPa to 5.0 GPa, and MX-type carbides each having a
5 circle-equivalent diameter of 10 nm or more exist with an average grain
spacing of 300 nm or less therebetween.
[0038] [2]
The steel material according to [1] contains, by mass%, one or two
selected from a group consisting of Cr: 0.05% to 0.25%, and Mo: 0.1%) to
10 0.35%.
[Effect of the Invention]
[0039] According to the present invention, it becomes possible to obtain
an impact absorbing member capable of suppressing or eliminating an
occurrence of crack thereon when an impact load is applied, and having a
15 high effective flow stress, so that it becomes possible to dramatically increase
an impact absorption energy of the impact absorbing member. By applying
the impact absorbing member as above, it becomes possible to further
improve a collision safety of a product of an automobile and the like, which is
industrially extremely useful.
20 [Brief Description of the Drawings]
[0040] [FIG. 1] FIG. 1 illustrates a heat pattern in continuous annealing
heat treatment employed in an example.
[Mode for Carrying out the Invention]
[0041] Hereinafter, the present invention will be described in detail. In
25 the following description, % related to a chemical composition of steel
indicates mass%.
10
[0042] 1. Chemical composition
Note that "%" in the following description regarding the chemical
composition means "mass%", unless otherwise noted.
[0043] (1) C: greater than 0.05% to 0.18%
5 C has a function of facilitating a generation of bainite being a main
phase, and austenite being a second phase, a function of improving a yield
strength and a tensile strength by increasing a strength of the second phase,
and a function of improving the yield strength and the tensile strength by
strengthening a steel through solid-solution strengthening. Further, C has a
10 function of coupling with Ti and V to precipitate MX-type fine carbides, and
improving the yield strength and a work hardening coefficient in a low-strain
region. If a C content is 0.05% or less, it is sometimes difficult to achieve an
effect provided by the above-described functions. Therefore, the C content
is set to be greater than 0.05%. On the other hand, if the C content exceeds
15 0.18%, there is a case where martensite and austenite are excessively
generated, which sometimes facilitates the occurrence of crack at the time of
applying the impact load. Therefore, the C content is set to 0.18% or less.
The C content is preferably 0.15% or less, and is more preferably 0.13%) or
less. Note that the present invention includes where the C content is
20 0.18%.
[0044] (2) Mn: 1% to 3%
Mn has a function of facilitating a generation of bainite by increasing
a hardenability, and a function of improving the yield strength and the tensile
strength by strengthening the steel through solid-solution strengthening. If a
25 Mn content is less than 1%, it is sometimes difficult to achieve an effect
provided by the above-described functions. Therefore, the Mn content is set
11
to 1% or more. The Mn content is preferably 1.5% or more. On the other
hand, if the Mn content exceeds 3%, there is a case where martensite and
austenite are excessively generated, resulting in that the local ductility is
significantly lowered. Therefore, the Mn content is set to 3% or less. The
5 Mn content is preferably 2.5% or less. Note that the present invention
includes a case where the Mn content is 1% and a case where the Mn content
is 3%.
[0045] (3) Si: greater than 0.5% to 1.8%
Si has a function of improving a uniform ductility and the local
10 ductility by suppressing a generation of carbide in bainite and martensite, and
a function of improving the yield strength and the tensile strength by
strengthening the steel through solid-solution strengthening. If a Si content
is 0.5% or less, it is sometimes difficult to achieve an effect provided by the
above-described functions. Therefore, the Si amount is set to be greater than
15 0.5%. The Si amount is preferably 0.8% or more, and is more preferably 1%
or more. On the other hand, if the Si content exceeds 1.8%, there is a case
where austenite excessively remains, and the impact crack sensitivity
becomes significantly high. Therefore, the Si content is set to 1.8% or less.
The Si content is preferably 1.5% or less, and is more preferably 1.3% or less.
20 Note that the present invention includes a case where the Si content is 1.8%.
[0046] (4) Al: 0.01% to 0.5%
Al has a function of suppressing a generation of inclusion in a steel
through deoxidation, and preventing the impact crack. If an Al content is
less than 0.01%, it is difficult to achieve an effect provided by the
25 above-described function. Therefore, the Al content is set to 0.01% or more.
On the other hand, if the Al content exceeds 0.5%, an oxide and a nitride
12
become coarse, which facilitates the impact crack, instead of preventing the
impact crack. Therefore, the Al content is set to 0.5% or less. Note that the
present invention includes a case where the Al content is 0.01% and a case
where the Al content is 0.5%.
5 [0047] (5) N: 0.001% to 0.015%
N has a function of suppressing a grain growth of austenite and ferrite
by generating a nitride, and suppressing the impact crack by refining a
structure. If a N content is less than 0.001%o, it is difficult to achieve an
effect provided by the above-described function. Therefore, the N content is
10 set to 0.001% or more. On the other hand, if the N content exceeds 0.015%o,
a nitride becomes coarse, which facilitates the impact crack, instead of
suppressing the impact crack. Therefore, the N content is set to 0.015%) or
less. Note that the present invention includes a case where the N content is
0.001% and a case where the N content is 0.015%.
15 [0048] (6) One or both of V and Ti: 0.01% to 0.3% in total
V and Ti have a function of generating carbides such as VC and TiC in
the steel, suppressing a growth of coarse crystal grains through a pinning
effect with respect to a grain growth of. ferrite, and suppressing the impact
crack. Further, V and Ti have a function of improving the yield strength and
20 the tensile strength by strengthening •' the steel through precipitation
strengthening realized by VC and TiC. Therefore, one or both of V and Ti is
(are) contained. If a total content of V and Ti (also referred to as "(V + Ti)
content", hereinafter) is less than 0.01%, it is difficult to achieve an effect
provided by the above-described functions. Therefore, the (V + Ti) content
25 is set to 0.01% or more. On the other hand, if the (V + Ti) content exceeds
0.3%o, VC or TiC is excessively generated, which increases the impact crack
13
sensitivity, instead of lowering the impact crack sensitivity. Therefore, the
(V + Ti) content is set to 0.3% or less. The present invention includes a case
where the total content of V and Ti is 0.01% and a case where the total
content is 0.3%. Any one of a case where only V is contained in an amount
5 of 0.01% to 0.3%, a case where only Ti is contained in an amount of 0.01% to
0.3%o, and a case where both of V and Ti are contained in an amount of 0.01%
to 0.3% in total, may be employed.
[0049] Further, it is also possible that one or two of Cr and Mo is (are)
contained as an optionally contained element.
10 [0050] (7) Cr: 0% to 0.25%
Cr is an optionally contained element, and has a function of increasing
a hardenability to facilitate a generation of bainite, and a function of
improving the yield strength and the tensile strength by strengthening the steel
through solid-solution strengthening- In order to more securely achieve
15 these functions, a content of Cr is preferably 0.05% or more. However, if
the Cr content exceeds 0.25%, a martensite phase is excessively generated,
which increases the impact crack sensitivity. Therefore, the Cr content is set
to 0.25%o or less. Note that the present invention includes a case where the
content of Cr is 0.25%.
20 [0051] . (8) Mo: 0% to 0.35%
Mo is, similar to Cr, an optionally contained element, and has a
function of increasing the hardenability to facilitate a generation of bainite
and martensite, and a function of improving the yield strength and the tensile
strength by strengthening the steel through solid-solution strengthening. In
25 order to more securely achieve these functions, a content of Mo is preferably
0.1% or more. However, if the Mo content exceeds 0.35%, the martensite
14
phase is excessively generated, which increases the impact crack sensitivity.
Therefore, when Mo is contained, the content of Mo is set to 0.35% or less.
Note that the present invention includes a case where the content of Mo is
0.35%.
5 [0052] The steel material of the present invention contains the
above-described essential contained elements, further contains the optionally
contained elements according to need, and contains a balance composed of Fe
and impurities. As the impurity, one contained in a raw material of ore,
scrap and the like, and one contained in a manufacturing step can be
10 exemplified. However, it is allowable that the other components are
contained within a range in which the properties of steel material intended to
be obtained in the present invention are not inhibited. For example,
although P and S are contained in the steel as impurities, P and S are desirably
limited in the following manner.
15 [0053] P: 0.02% or less
P makes a grain boundary to be fragile, and deteriorates a hot
workability. Therefore, an upper limit of P content is set to 0.02% or less.
It is desirable that the P content is as small as possible, but, based on the
assumption that a dephosphorization is performed within a range of actual
20 manufacturing steps and manufacturing cost, the upper limit of P content is
0.02%). The upper limit is desirably 0.015% or less.
[0054] S: 0.005%o or less
S makes the grain boundary to be fragile, and deteriorates the hot
workability and ductility. Therefore, an upper limit of P content is set to
25 0.005%o or less. It is desirable that the S content is as small as possible, but,
based on the assumption that a desulfurization is performed within a range of
15
actual manufacturing steps and manufacturing cost, the upper limit of S
content is 0.005%. The upper limit is desirably 0.002% or less.
[0055] 2. Steel structure
A steel structure related to the present invention contains bainite with
5 fine block size as a main phase, and further, it improves the plastic flow stress
with the use of fine precipitates, in order to realize both of an increase in
effective flow stress by obtaining a high yield strength and a high work
hardening coefficient in the low-strain region, and an impact crack resistance.
[0056] (1) Area ratio of bainite: 80% or more
10 If an area ratio of bainite being the main phase is less than 80%, it
becomes difficult to secure a high yield strength. Therefore, the area ratio of
bainite being the main phase is set to 80% or more. The area ratio of bainite
is preferably 85% or more, and is more preferably greater than 90%.
[0057] (2) Average block size of bainite: less than 2.0 jum
15 The ductility can be increased by refining bainite being the main
phase. If an average block size of bainite is 2.0 jam or more, it is difficult to
improve the ductility. Therefore, the average block size of bainite is set to
less than 2.0 jim. This block size is preferably 1.5 (im or less.
[0058] (3) One or two or more selected from a group consisting of ferrite,
20 martensite and austenite is (are) contained in an amount of 5% or more in
total, and an average grain diameter of all of the above-described ferrite,
martensite and bainite is less than 1.0 urn.
If it is set that in the steel material containing bainite as the main
phase, a second phase thereof contains one or two or more selected from a
25 group consisting of ferrite, martensite and austenite, and these elements are
refined, the local ductility can be further improved. If a total area ratio of
16
ferrite, martensite and austenite is less than 5%, or if an average grain
diameter of all of ferrite, martensite and austenite is 1.0 um or more, it is
difficult to further improve the local ductility. Therefore, it is set that one or
two or more selected from a group consisting of ferrite, martensite and
5 austenite is (are) contained in an amount of 5% or more in total, and the
average grain diameter of all of the above-described ferrite, martensite and
austenite is less than 1.0 um.
[0059] Note that if ferrite is contained in the second phase, the fracture
toughness can be improved, if austenite is contained in the second phase, the
10 uniform elongation can be improved, and if martensite is contained in the
second phase, the strength can be increased. There is a case where, other
than ferrite, martensite and austenite, cementite and perlite are inevitably
contained in the second phase other than bainite being the main phase, and
such an inevitable structure is allowed to be contained if the structure is 5
15 area% or less.
[0060] (4) Average nanohardness of bainite: not less than 4.0 GPa nor
more than 5.0 GPa
If an average nanohardness of bainite is less than 4.0 GPa, it becomes
difficult to secure a tensile strength of 980 MPa or more in a steel material in
20 which the area ratio of bainite is 80% or more. Therefore, the average
nanohardness of bainite is set to 4.0 GPa or more. On the other hand, if the
average nanohardness of bainite exceeds 5.0 GPa, it becomes difficult to
suppress the occurrence of crack when applying the impact load. Therefore,
the average nanohardness of bainite is set to 5.0 GPa or less.
25 [0061] Here, the nanohardness is a value obtained by measuring a
nanohardness in a bainite block by using a nanoindentation. In the present
17
invention, a cube corner indenter is used, and a nanohardness obtained under
an indentation load of 500 u,N is adopted.
[0062] (5) Average grain spacing of MX-type carbides each having
circle-equivalent diameter of 10 nm or more: 300 nm or less
5 In the steel material containing bainite as the main phase, a
precipitation site of the second phase is a prior austenite grain boundary, and
in order to refine the second phase, it is necessary to refine austenite grains.
As a result of studying various methods for refining austenite grains, it was
clarified that by employing suitable hot-rolling conditions and heat treatment
10 conditions to obtain a pinning effect provided by MX-type carbides, a growth
of coarse crystal grains can be greatly suppressed, as will be described later.
[0063] The MX-type carbide is a carbide having a NaCl-type crystal
structure, and is formed of V and/or Ti and C. A size of the MX-type
carbide exhibiting the pinning effect is 10 nm or more in a circle-equivalent
15 diameter. If the size of the MX-type carbide is less than 10 nm in the
circle-equivalent diameter, the pining effect with respect to a grain boundary
migration cannot be expected. Therefore, the refining of structure is tried to
be realized by making the MX-type carbides each having the circle-equivalent
diameter of 10 nm or more exist, but, if an average grain spacing between the
20 carbides exceeds 300 nm, it is difficult to achieve a sufficient pinning effect
Therefore, it is set that the MX-type carbides each having the
circle-equivalent diameter of 10 nm or more exist with the average grain
spacing of 300 nm or less therebetween.
[0064] A density of the MX-type carbides each having the
25 circle-equivalent diameter of 10 nm or more is preferably as high as possible,
so that a lower limit of the average grain spacing between the carbides is not
18
particularly specified, but, realistically, the lower limit is 50 nm or more.
Although an upper limit of the size of the MX carbide is not particularly
specified, an excessively coarse size may exert an adverse effect on the
ductility, instead of improving the ductility, so that the upper limit of the size
5 of the MX carbide (circle-equivalent diameter) is preferably set to 50 nm.
[0065] 3. Properties
The steel material according to the present invention has a
characteristic in a point that the effective flow stress is high, the impact
absorption energy is high, and at the same time, the occurrence of crack when
10 applying the impact load is suppressed. This characteristic is proved based
on a high 5% flow stress, a high average crush load, and a high stable
buckling ratio in a buckling test, as will be indicated in later-described
examples. The 5% flow stress is preferably 700 MPa or more.
[0066] As other mechanical properties, there can be cited properties in
15 which the strength is high and the ductility and a hole expandability are
excellent, such that the tensile strength is 982 MPa or more, the uniform
elongation (total elongation) is 7% or more, and a hole expansion ratio is
120% or more when measured by a measurement method based on Japan Iron
and Steel Federation standard JFST 1001-1996.
20 [0067] . 4. Manufacturing method .
The steel material of the present invention can be obtained through the
following manufacturing methods (1) to (3), for example.
[0068] Manufacturing method (1): hot-rolled material (no performance of
heat treatment)
25 In order to obtain the steel material of the present invention as
hot-rolled, it is preferable to properly precipitate VC and TiC in a hot-rolling
19
step to suppress a growth of coarse crystal grains with the use of the pinning
effect provided by VC and TiC, and to optimize a multi-phase structure by
controlling a thermal history.
[0069] First, a slab having the above-described chemical composition is
5 set to have a temperature of 1200°C or more and subjected to multi-pass
rolling at a total reduction ratio of 50% or more, and the rolling is completed
in a temperature region of not less than 800°C nor more than 950°C. Within
a period of time of 0.4 seconds after the completion of the rolling, the
resultant is cooled at a cooling rate of 600°C/second or more to a temperature
10 region of 500°C or less, and coiled in a temperature region of not less than
300°C nor more than 500°C, to thereby produce a hot-rolled steel sheet.
[0070] Through the above-described hot rolling and cooling, it is possible
to obtain a steel structure as hot-rolled, having the MX-type carbides
dispersed therein, and mainly formed of a bainite structure with a fine block
15 size.
[0071] When the above-described hot-rolling conditions are not satisfied,
there is a case where an intended steel structure cannot be obtained and the
ductility and the strength are lowered, since austenite becomes coarse, and
besides, a precipitation density of the MX-type carbides is decreased.
20 Further, when the above-described cooling conditions are not satisfied, there
is a case where the generation of ferrite in the cooling step becomes excessive,
and besides, the block size of bainite becomes too large, resulting in that
desired impact properties cannot be achieved.
[0072] In this manufacturing method (1), after the hot rolling is
25 practically completed, rapid cooling is conducted at a cooling rate of
600°C/second or more to a temperature region of 500°C or less within a
20
period of time of 0.4 seconds. The practical completion of hot rolling means
a pass in which the practical rolling is conducted at last, in the rolling of
plurality of passes conducted in finish rolling of the hot rolling. For
example, in a case where the practical final reduction is conducted in a pass
5 on an upstream side of a finishing mill, and the practical rolling is not
conducted in a pass on a downstream side of the finishing mill, the rapid
cooling is conducted to the temperature region of 500°C or less within a
period of time of 0.4 seconds after the rolling in the pass on the upstream side
is completed. Further, for example, in a case where the practical rolling is
10 conducted up to when the pass reaches the pass on the downstream side of the
finishing mill, the rapid cooling is conducted to the temperature region of
500°C or less within a period of time of 0.4 seconds after the rolling in the
pass on the downstream side is completed. Note that the rapid cooling is
basically conducted by a cooling nozzle disposed on a run-out-table, but, it is
15 also possible to be conducted by an inter-stand cooling nozzle disposed
between the respective passes of the finishing mill.
[0073] The above-described cooling rate (600°C/second or more) is set
based on a temperature of a surface of sample (surface temperature of steel
sheet) measured by a thermotracer. A cooling rate (average cooling rate) of
20 the entire, steel sheet is estimated to be about 200°C/second or more, as a
result of conversion from the cooling rate (600°C/second or more) based on
the surface temperature.
[0074] Manufacturing method (2): Hot-rolled and heat-treated material
In order to obtain the steel material of the present invention by
25 performing heat treatment after hot rolling, it is preferable that VC and TiC
are properly precipitated in a hot-rolling step and a temperature-raising
21
process in a heat treatment step, a growth of coarse crystal grains is
suppressed by a pinning effect provided by VC and TiC, and an optimization
of multi-phase structure is realized during the heat treatment.
[0075] First, a slab having the above-described chemical composition is
5 set to have a temperature of 1200°C or more and subjected to multi-pass
rolling at a total reduction ratio of 50% or more, and the rolling is completed
in a temperature region of not less than 800°C nor more than 950°C. Within
a period of time of 0.4 seconds after the completion of the rolling, the
resultant is cooled at a cooling rate of 600°C/second or more to a temperature
10 region of 700°C or less (this cooling is also referred to as primary cooling),
and then cooled to a temperature region of 500°C or less at a cooling rate of
less than 100°C/second (this cooling is also referred to as secondary cooling),
and after that, the resultant is coiled in a temperature region of not less than
300°C nor more than 500°C, to thereby produce a hot-rolled steel sheet.
15 [0076] By this hot-rolling step, the hot-rolled steel sheet in which the
MX-type carbides are precipitated at high density in the ferrite grain boundary,
is obtained. On the other hand, when the above-described hot-rolling
conditions are not satisfied, it becomes difficult to obtain the steel material of
the present invention since the average grain diameter of the MX-type
20 carbides becomes too small and the pinning effect with respect to the grain
growth is reduced, and an average intergranular distance of the MX-type
carbides becomes too large, which does not contribute to the refining of
crystal grains.
[0077] In this manufacturing method (2), after the hot rolling is
2.5 practically completed, rapid cooling is conducted at a cooling rate of
600°C/second or more to a temperature region of 700°C or less within a
22
period of time of 0.4 seconds. Similar to the previously described
manufacturing method (1), also in the manufacturing method (2), the practical
completion of hot rolling means a pass in which the practical rolling is
conducted at last, in the rolling of plurality of passes conducted in finish
5 rolling of the hot rolling. The rapid cooling is basically conducted by a
cooling nozzle disposed on a run-out-table, but, it is also possible to be
conducted by an inter-stand cooling nozzle disposed between the respective
passes of the finishing mill.
[0078] The above-described cooling rate (600°C/second or more) is set
10 based on a temperature of a surface of sample (surface temperature of steel
sheet) measured by a thermotracer. A cooling rate (average cooling rate) of
the entire steel sheet is estimated to be about 200°C/second or more, as a
result of conversion from the cooling rate (600°C/second or more) based on
the surface temperature.
15 [0079] In this manufacturing method (2), next, a temperature of the
hot-rolled steel sheet obtained by the above-described hot-rolling step is
raised to a temperature region of not less than 850°C nor more than 920°C at
an average temperature rising rate of not. less than 2°C/second nor more than
50°C/second, and the steel sheet is retained in the temperature region for a
20 period of time of not less than 100 seconds nor more than 300 seconds
(annealing in FIG. 1). Subsequently, heat treatment in which the resultant is
cooled to a temperature region of not less than 270°C nor more than 390°C at
an average cooling rate of not less than 10°C/second nor more than
50°C/second, and retained in the temperature region for a period of time of
25 not less than 10 seconds nor more than 300 seconds, is performed (quenching
in FIG 1).
I'
23
[0080] If the above-described average temperature rising rate is less than
2°C/second, the grain growth of ferrite occurs during the temperature rising,
resulting in that the crystal grains become coarse. Although the
above-described average temperature rising rate is preferably as high as
5 possible, realistically, it is 50°C/second or less. If the temperature retained
after the above-described temperature rising is less than 850°C or the
retention time is less than 100 seconds, an austenitize required for the
quenching becomes insufficient, resulting in that it becomes difficult to obtain
an intended multi-phase structure. On the other hand, if the temperature
10 retained after the above-described temperature rising exceeds 920°C or the
retention time exceeds 300 seconds, austenite becomes coarse, resulting in
that it becomes difficult to obtain an intended multi-phase structure.
[0081] After the above-described temperature rising, in order to obtain a
structure mainly formed of bainite, it is necessary to perform quenching at a
15 bainite transformation temperature or less while suppressing a ferrite
transformation. If the above-described average cooling rate is less than
10°C/second, a ferrite amount becomes excessive, and it is difficult to obtain
a sufficient strength. Although the above-described average cooling rate is
preferably as high as possible, realistically, it is 50°C/second or less. Further,
20 if a cooling stop temperature of the cooling described above is less than
270°C, an area ratio of martensite becomes too large, resulting in that the
local ductility is lowered. On the other hand, if the cooling stop temperature
of the cooling described above exceeds 390°C, the average block size of
bainite becomes coarse, resulting in that the strength and the ductility are
25 lowered. Further, if the retention time in the temperature region of not less
than 270°C nor more than 390°C is less than 10 seconds, the facilitation of
24
bainite transformation sometimes becomes insufficient. On the other hand,
if the retention time in the temperature region of not less than 270°C nor more
than 390°C exceeds 300 seconds, the productivity is significantly hindered.
[0082] It is also possible to adjust a hardness of bainite by conducting,
5 after the above-described quenching, tempering treatment according to need
in which a retention is performed in a temperature region of not less than
400°C nor more than 550°C for a period of time of not less than 10 seconds
nor more than 650 seconds (tempering 1 and tempering 2 in FIG. 1). Note
that the tempering may be performed in one stage, or may also be performed
10 in a plurality of stages separately. FIG. 1 illustrates an example in which the
tempering is performed in two stages separately.
[0083] Here, if the tempering temperature is less than 400°C or the
tempering time is less than 10 seconds, it is not possible to sufficiently
achieve an effect provided by the tempering. On the other hand, if the
15 tempering temperature exceeds 550°C or the tempering time exceeds 650
seconds, there is a case where an intended strength cannot be obtained due to
the decrease in strength. The tempering can be conducted through heating in
two stages or more within the above-described temperature region. In that
case, it is preferable that a heating temperature in the first stage is set to be
20 lower than a heating temperature in the second stage.
[0084] Manufacturing method (3): Cold-rolled and heat-treated material
In order to obtain the steel material of the present invention by
performing heat treatment after hot rolling and cold rolling, it is preferable
that VC and TiC are properly precipitated in a hot-rolling step and a
25 temperature-raising process in a heat treatment step, a growth of coarse
crystal grains is suppressed by a pinning effect provided by VC and TiC, and
25
an optimization of multi-phase structure is realized during the heat treatment,
similar to the manufacturing method (2). In order to achieve the above, it is
preferable to perform manufacture through a manufacturing method including
the following steps.
5 [0085] First, a slab having the above-described chemical composition is
set to have a temperature of 1200°C or more and subjected to multi-pass
rolling at a total reduction ratio of 50% or more, and the rolling is completed
in a temperature region of not less than 800°C nor more than 950°C. Within
a period of time of 0.4 seconds after the completion of the rolling, the
10 resultant is cooled at a cooling rate of 600°C/second or more to a temperature
region of 700°C or less (this cooling is also referred to as primary cooling),
and then cooled to a temperature region of 500°C or less at a cooling rate of
less than 100°C/second (this cooling is also referred to as secondary cooling),
and after that, the resultant is coiled in a temperature region of not less than
15 300°C nor more than 500°C, to thereby produce a hot-rolled steel sheet.
[0086] By this hot-rolling step, the hot-rolled steel sheet in which the
MX-type carbides are precipitated at high density in the ferrite grain boundary,
is obtained. On the other hand, when the above-described hot-rolling
conditions are not satisfied, it becomes difficult to obtain the steel material of
20 the present invention since the average grain diameter of the MX-type
carbides becomes too small and the pinning effect with respect to the grain
growth is reduced, and an average intergranular distance of the MX-type
carbides becomes too large, which does not contribute to the refining of
crystal grains.
25 [0087] In this manufacturing method (3), after the hot rolling is
practically completed, rapid cooling is conducted at a cooling rate of
26
600°C/second or more to a temperature region of 700°C or less within a
period of time of 0.4 seconds. Similar to the previously described
manufacturing methods (1) and (2), also in the manufacturing method (3), the
practical completion of hot rolling means a pass in which the practical rolling
5 is conducted at last, in the rolling of plurality of passes conducted in finish
rolling of the hot rolling. The rapid cooling is basically conducted by a
cooling nozzle disposed on a run-out-table, but, it is also possible to be
conducted by an inter-stand cooling nozzle disposed between the respective
passes of the finishing mill.
10 [0088] The above-described cooling rate (60Q°C/second or more) is set
based on a temperature of a surface of sample (surface temperature of steel
sheet) measured by a thermotracer. A cooling rate (average cooling rate) of
the entire steel sheet is estimated to be about 200°C/second or more, as a
result of conversion from the cooling rate (600°C/second or more) based on
15 the surface temperature.
[0089] In this manufacturing method (3), next, cold rolling at a reduction
ratio of not less than 30% nor more than 70% is conducted to produce a
cold-rolled steel sheet.
[0090] Next, a temperature of the cold-rolled steel sheet obtained by the
20 above-described cold-rolling step is raised to a temperature region of not less
than 850°C nor more than 920°C at an average temperature rising rate of not
less than 2°C/second nor more than 50°C/second, and the steel sheet is
retained in the temperature region for a period of time of not less than 100
seconds nor more than 300 seconds (annealing in FIG. 1). Subsequently,
25 heat treatment in which the resultant is cooled to a temperature region of not
less than 270°C nor more than 390°C at an average cooling rate of not less
27
than 10°C/second nor more than 50°C/second, and retained in the temperature
region for a period of time of not less than 10 seconds nor more than 300
seconds, is performed (quenching in FIG. 1).
[0091] If the above-described average temperature rising rate is less than
5 2°C/second, the grain growth of ferrite occurs during the temperature rising,
resulting in that the crystal grains become coarse. Although the
above-described average temperature rising rate is preferably as high as
possible, realistically, it is 50°C/second or less. If the temperature retained
after the above-described temperature rising is less than 850°C or the
10 retention time is less than 100 seconds, an austenitize required for the
quenching becomes insufficient, resulting in that it becomes difficult to obtain
an intended multi-phase structure. On the other hand, if the temperature
retained after the above-described temperature rising exceeds 920°C or the
retention time exceeds 300 seconds, austenite becomes coarse, resulting in
15 that it becomes difficult to obtain an intended multi-phase structure.
[0092] After the above-described temperature rising, in order to obtain a
structure mainly formed of bainite, it is necessary to perform quenching at a
bainite transformation temperature or less while suppressing a ferrite
transformation. If the above-described average cooling rate is less than
20 10°C/second, a ferrite amount becomes excessive, and it is difficult to obtain
a sufficient strength. Although the above-described average cooling rate is
preferably as high as possible, realistically, it is 50°C/second or less. Further,
if a cooling stop temperature of the cooling described above is less than
270°C, an area ratio of martensite becomes too large, resulting in that the
25 local ductility is lowered. On the other hand, if the cooling stop temperature
of the cooling described above exceeds 390°C, the average block size of
28
bainite becomes coarse, resulting in that the strength and the ductility are
lowered. Further, if the retention time in the temperature region of not less
than 270°C nor more than 390°C is less than 10 seconds, the facilitation of
bainite transformation sometimes becomes insufficient. On the other hand,
5 if the retention time in the temperature region of not less than 270°C nor more
than 390°C exceeds 300 seconds, the productivity is significantly hindered.
[0093] It is also possible to adjust a hardness of bainite by conducting,
after the above-described quenching, tempering treatment according to need
in which a retention is performed in a temperature region of not less than
10 400°C nor more than 550°C for a period of time of not less than 10 seconds
nor more than 650 seconds, similar to the previously described manufacturing
method (2). Here, if the tempering temperature is less than 400°C or the
tempering time is less than 10 seconds, it is not possible to sufficiently
achieve an effect provided by the tempering. On the other hand, if the
15 tempering temperature exceeds 550°C or the tempering time exceeds 650
seconds, there is a case where an intended strength cannot be obtained due to
the decrease in strength. The tempering can be conducted through heating in
two stages or more within the above-described temperature region. In that
case, it is preferable that a heating temperature in the first stage is set to be
20 lower than a heating temperature in the second stage.
[0094] The hot-rolled steel sheet or the cold-rolled steel sheet
manufactured through the manufacturing methods (1) to (3) as above may be
used as it is as the steel material of the present invention, or a steel sheet, cut
from the hot-rolled steel sheet or the cold-rolled steel sheet, on which
25 appropriate working such as bending and presswork is performed according to
need, may also be employed as the steel material of the present invention.
29
Further, the steel material of the present invention may also be the steel sheet
as it is, or the steel sheet on which plating is performed after the working.
The plating may be either electroplating or hot dipping, and although there is
no limitation in a type of plating, the type of plating is normally zinc or zinc
5 alloy plating.
[Examples]
[0095] An experiment was conducted by using slabs (each having a
thickness of 35 mm, a width of 160 to 250 mm, and a length of 70 to 140
mm) having chemical compositions presented in Table 1. In Table 1, "-"
10 means that the element is not contained positively. An underline indicates
that a value is out of the range of the present invention. A steel type D is a
comparative example in which a total content of V and Ti is less than the
lower limit value. A steel type I is a comparative example in which a
content of Mn exceeds the upper limit value. A steel type J is a comparative
15 example in which a content of C exceeds the upper limit value. In each of
the steel types, a molten steel of 150 kg was produced in vacuum to be cast,
the resultant was then heated at a furnace temperature of 1250°C, and
subjected to hot forging at a temperature of 950°C or more, to thereby obtain
a slab.
30
[0096]
[Table 1]
TYPE
A
B
C
D
E
F
G
H
I
J
CHEMICAL COMPOSITION (UNIT: MASSES, BALANCE: Fe AND IMPURITIES)
C
0.12
0.12
0.12
0.12
Q.12
0.18
0.1.5
0.1S
0.15
0.22
Si
1.24
1.23
1.25
1.23
1.4S
1.25
1.30
1.33
1.52
1.32
Mn
2,05
2,01
2.01
2,25
2,02
2,20
2,02
2.20
12
2.15
P
o.oos
0.009
0.00?
0.0 SI
0.013
0.010
0.012
0.010
0.012
0.010
S
0.002
0.002
0.002
0.002
0,003
0.003
0.002
0.002
0.002
O.O02
Cr
0.12
0.20
0.15
0 10
0.10
-
0.10
0,10
0.15
0.15
Mo
-
0.2O
-
-
.
-
-
0.22
-
-
V
0.20
0.15
0,05
;
0.25
0.20
0.25
-
0.20
-
Ti
0.005
0.005
0.005
;
0.005
0.003
-
0.012
0.004
0.005
Al
0.033
0.030
•0.032
0.035
0.033
0.051
0.035
0.35
0.035
0.025
N
0.0024
0.0025
0.0026
0.0045
0.0025
13.0031
0,0024
0,0025
0 0035
0.0032
UNDERLINE INDICATES THAT VALUE IS OUT OF RANGE OF PRESENT INVENTION
[0097] Each of the above-described slabs was reheated at 1250°C within
1 hour, and after that, the resultant was subjected to rough hot rolling in 4
5 passes by using a hot-rolling testing machine, the resultant was further
subjected to finish hot rolling in 3 passes, and after the completion of rolling,
primary cooling and secondary cooling were conducted, to thereby obtain a
hot-rolled steel sheet. Hot-rolling conditions are presented in Table 2, The
primary cooling and the secondary cooling right after the completion of
10 rolling were conducted by water cooling. The secondary cooling was
completed at a coiling temperature presented in Table.
15
-
i
i
2
3
4
J
5
7
i
9
10
11
12
15
14
IS
IS
1?
I
A
A
A
A
A
B
c
a
E
E
E
E
F
G
H
I
i
HOT ROLLINCROUCH
ROLLING
TOTAL
REDUCTION
RATIO
33
Si
S3
S3
S3
S3
S3
S3
S3
S3
S3
S3
33
S3
S3
S3
S3
FINISH HOT ROLLING
NUMBER
OF
PASSES
REDUCTION
RATIO
IN EACH PASS
30°v303i-J0^i
50°v30^-30%
3G?o-30%-39*i
30°i-30c-&-3«"e
>0%-tO%-l,
3f!%.}O3i-30?'.
30%-J0%-301€00
?1000
>1000
>1000
>1'MC
MCsOC
>i000
>1000
>1QOO
M0Q0
--100C
COOLINGSTOP
•TEMPERATURE
Co;
450
-4:0
55C
553
fc"C
450
650
6:0
630
6:0
630
630
?IO00 - 650
>im
?lO0C
>iooc
>iooo
650
650
650
650
?EKiCJ>OF
0= ROLLING
TO START
OT COOUNG
CM
IS
0.1
0.1
Li
0.1
0.1
0.1
0.1
0.1
0.1
0.1
0.1
SECONDARY
COOLING
AVER-AGE
COOLING
RATE
r-c.s)
—
-
17
15
10
-
I?
IS
1?
15
IS
16
19
0.1 19
0,1
0.1
0.1
19
16
19
COOLING
STOP
TEMPERATURE
C-Ci
-
-
415
460
450
__
4 1 7
420
420
455
400
455
430
450
410
460
410
COILI
TEMPERA
450
450
400
450
450
45
40
40
40
45
45
45
40
40
40
42
40
UNDERLINE IKDICATE5 THAT VALUE IS OUT OF RANGE OF PRESENT IN~'ENTION
32
[0099] The steel sheets of test numbers 1, 2, 6, 13, and 15 to 17 were set
to be steel sheets as hot-rolled, without performing cold rolling. On the
other steel sheets of test numbers 3 to 5, 7 to 12, and 14, the cold rolling was
performed. As can be understood from Table 2 and Table 3, a sheet
5 thickness of each of the obtained hot-rolled steel sheets or cold-rolled steel
sheets was 1.6 mm. On the steel sheets of test numbers 4, 5, 9 to 12, and 14,
heat treatment was performed by using a continuous annealing simulator with
a heat pattern presented in FIG. 1 and under conditions presented in Table 3.
In the present examples, a process from a temperature rising to a temperature
10 retention in the heat treatment corresponds to annealing, cooling after the
annealing corresponds to quenching, and heat treatment thereafter
corresponds to tempering conducted for the puipose of performing hardness
adjustment (softening). As can be understood from FIG. 1 and Table 3, the
tempering heat treatment in the temperature region of not less than 400°C nor
15 more than 550°C was conducted in two stages. Note that on the steel sheets
of test numbers 3, 7, 8, and 13, only the quenching was performed after the
annealing, and the tempering was not performed.
-z
I
:
?
4
5
6
-
S
9
10
11
12
13
14
15
Is
i :
d
A
A
A
A
A
3
C
D
E
Z
E
E
F
fT
H
I
J
TOTAL
REDUCTION
RATIO
IN COLD
ROLLING
AS HOT-ROLLED
AS HOT-ROLLED
50%
:QH
50%
AS HOT-ROLLED
fO'c
50?;-
50%
50%
50?-7
50%
AS HOT-ROLLED
50? i
AS HOT-ROLLED
AS HOT-ROLLED
AS HOT-ROLLED
CONDITIONS OF CONTINUOUS ANNEALING
CONDITIONS
OF ANNEALING
TEMPERATURE
RISING
RATE
f'C's'i
--
10
10
10
-
10
10
10
10
10
ANNEAL ENG
TEMPERATURE
•CO
--
900
9O0
900
-
92t
92C
90S
850
S50
20 I 900
10
10
---
S5&
900
-
-
-
ANNTALING
TEME
_
-
250
2:0
2:0
_
250
' 250
250
250
120
120
250
250
_
--
CGQLLNG
RATE
-
_
40
41}
40
-
•5
55
40
40
40
5
40
40 •
-
_
_
CONDITIONS FROM QUENCHING
TO TEMPESING < TO ©)
QUENCHING
TE>.iPERATUXE
CO
QUENCHING
TIME
j
-
350
530
" 330
-
• to
330
330
530
2^
530
•30
330
---
-
120
120
120
-
120
L:O
I:O
120
s®
120
120
120
-
_
-
TEMPERING
TEMPERATURE
0>
CO
-
_
_
460
460
--
_
460
460
4S0
4(50
_
4S0
-
_
_
TEMPERING
TD-fE

II
!2
13
14
15
Id
IT
S
£
A
A
A
A
A
3
C
o
E
E
E
E
F
G
H
I
;
STEEL STRUCTURE
AREA
EATIQ
OF
EAL\TTE
pi)
33
92
«
92
s;
91
9:
21
94
7~
AVESAC-E
BLOCK
SIZE OF
3AINITE
fern)
1.2
11
1.4
;.;
LI
I.I
1.2
M
l.i
U
o