[0001] The present invention relates to a high strength thin steel sheet that has high hydrogen
embrittlement resisting property (particularly the hydrogen embrittlement resisting
property after being subjected to forming process) and high workability, especially
to a high strength thin steel sheet that has high resistance against fractures due
to hydrogen embrittlement such as season crack and delayed fracture that pose serious
problems for steel sheets having tensile strength of 1180 MPa or higher, and has high
workability.
[0002] There are increasing demands for the steel sheet, that is pressed or bent into a
form of a high-strength component of automobile or industrial machine, to have both
high strength and high ductility at the same time. In recent years, there are increasing
needs for high strength steel sheets having strength of 1180 MPa or higher, as the
automobiles are being designed with less weight. A type of steel sheet that is regarded
as promising to satisfy these needs is TRIP (transformation induced plasticity) steel
sheet.
[0003] The TRIP steel sheet includes residual austenite structure and, when processed to
deform, undergoes considerable elongation due to induced transformation of the residual
austenite (residual γ) into martensite by the action of stress. Known examples of
the TRIP steel include TRIP type composite-structure steel (TPF steel) that consists
of polygonal ferrite as the matrix phase and residual austenite; TRIP type tempered
martensite steel (TAM steel) that consists of tempered martensite as the matrix phase
and residual austenite; and TRIP type bainitic steel (TBF steel) that consists of
bainitic ferrite as the matrix phase and residual austenite. Among these, the TBF
steel has long been known (described, for example, in NISSIN STEEL TECHNICAL REPORT,
No. 43, Dec. 1980, pp1-10), and has such advantages as the capability to readily provide
high strength due to the hard bainitic ferrite structure, and the capability to show
outstanding elongation because fine residual austenite grains can be easily formed
in the boundary of lath-shaped bainitic ferrite in the bainitic ferrite structure.
The TBF steel also has such an advantage related to manufacturing, that it can be
easily manufactured by a single heat treatment process (continuous annealing process
or plating process).
[0004] In the realm of high strength of 1180 MPa upward, however, the TRIP steel sheet is
known to suffer a newly emerging problem of delayed fracture caused by hydrogen embrittlement,
similarly to the conventional high strength steel. Delayed fracture refers to the
failure of high-strength steel under stress, that occurs as hydrogen originating in
corrosive environment or the atmosphere infiltrates and diffuses in microstructural
defects such as dislocation, void and grain boundary, and makes the steel brittle.
This results in decreases in ductility and toughness of the metallic material.
[0005] It has been well known that the high strength steel that is widely used in the manufacture
of PC steel wire and line pipe experiences hydrogen embrittlement (pickling embrittlement,
plating embrittlement, delayed fracture, etc.) caused by the infiltration of hydrogen
into the steel when tensile strength of the steel becomes 980 MPa or higher. Accordingly,
most of technologies of improving hydrogen embrittlement resisting property have been
developed aiming at steel members such as bolt. "New Development in Elucidation of
Delayed Fracture" (published by The Iron and Steel Institute of Japan in January,
1997), for example, describes that it is effective in improving the resistance against
delayed fracture to add element such as Cr, Mo or V that demonstrates resistance against
temper softening to the metal structure that is based on tempered martensite as the
major phase. This technology is intended to cause the delayed fracture to take place
within grains instead of in the grain boundaries, thereby to constrain the fracture
from occurring, by precipitating alloy carbide and making use thereof as the site
for trapping hydrogen.
[0006] Thin steel sheets having strength higher than 780 MPa have rarely been used for the
reason of workability and weldability. Also hydrogen embrittlement has rarely been
regarded as a problem for thin steel sheets where hydrogen that has infiltrated therein
is immediately released due to the small thickness. For these reasons, much efforts
have not been dedicated to counter the hydrogen embrittlement. In recent years, however,
higher strength is required of the reinforcement members such as bumper, impact beam
and seat rail, etc., in order to meet the requirement of weight reduction of the automobile
and to improve the collision safety. Automobile components that are shaped by pressing
or bending process such as pillar are also required to have higher strength. As a
result, there have been increasing demands for high strength steel sheet having strength
of 980 MPa or higher for the manufacture of these parts. This makes it necessary to
improve hydrogen embrittlement resisting property of the high strength steel sheet.
[0007] Use of the technology addressed to the bolt steel described above may be considered
for improving the hydrogen embrittlement resisting property of the high strength steel
sheet. However, in the case of "New Development in Elucidation of Delayed Fracture"
(published by The Iron and Steel Institute of Japan in January, 1997), for example,
0.4% or higher of C content and much alloy elements are contained, and therefore application
of this technology to a thin steel sheet compromises the workability required of the
thin steel sheet. The technology also has a drawback related to the manufacturing
process, since it takes several hours or longer period of heat treatment to cause
the alloy carbide to precipitate. Therefore, improvement of the hydrogen embrittlement
resisting property of a thin steel sheet requires it to develop a novel technology.
[0008] It is relatively easy to achieve a high strength with quench-hardened (tempered)
martensite steel that has been commonly used as a high-strength steel. However, improvement
of the workability without variability essentially requires it to provide a tempering
process which makes it necessary to strictly control the temperature and duration
of the process. This also sometimes increases the possibility of tempering embrittlement
and makes it difficult to reliably improve workability. Although there is a steel
of composite structure of martensite and ferrite or the like developed to improve
ductility, such a steel has a high notch sensitivity due to mixed presence of hard
phase and soft phase, thus making it difficult to achieve sufficient improvement of
hydrogen embrittlement resisting property.
[0009] Hydrogen-induced delayed fracture is believed to occur in such a steel that contains
martensite, because hydrogen is concentrated in grain boundaries of prior austenite
thereby to form voids or other defects that become the starting points of the fracture.
Common practice that has been employed to decrease the sensitivity to delayed fracture
is to diffuse fine grains of carbide or the like uniformly as the site for trapping
hydrogen, thereby to decrease the concentration of diffusive hydrogen. However, even
when a large number of carbide grains or the like are diffused as the trap site for
hydrogen, there is a limitation to the hydrogen trapping capability and delayed fracture
attributable to hydrogen cannot be fully suppressed.
[0010] Japanese Unexamined Patent Publication (Kokai) No. 11-293383 describes a technology
to improve the hydrogen embrittlement resisting property of steel sheet, where hydrogen-induced
defects can be suppressed by having oxides that include Ti and Mg exist as the main
components in the structure. However, this technology is intended for thick steel
sheets and, although consideration is given to delayed fracture after welding with
a large input heat, no consideration is given to the environment (for example, corrosive
environment, etc.) in which automobile parts manufactured by using thin steel sheets
are used.
[0011] Japanese Unexamined Patent Publication (Kokai) No. 2003-166035 describes that it
is made possible to improve the ductility and delayed fracture resistance after being
subjected to forming process, by controlling the mutual relationships between 1) the
form (standard deviation and mean grain size) in which oxide, sulfide, composite crystallization
product or composite precipitate of Mg is dispersed, 2) volumetric proportion of residual
austenite and 3) strength of the steel sheet. However, it is difficult to improve
the hydrogen embrittlement resisting property in such an environment as hydrogen is
generated through corrosion of the steel sheet simply through the trapping effect
achieved by controlling the form of precipitate.
[0012] It has been a common practice in the past to reduce the residual austenite that was
believed to have an adverse effect on the hydrogen embrittlement resisting property.
In recent years, however, the effect of residual austenite on the improvement of hydrogen
embrittlement resisting property has been recognized and accordingly much attention
has been paid to the TRIP steel that contains residual austenite.
[0013] Tomohiko HOJO et. al "Hydrogen Embrittlement of High Strength Low Alloy TRIP Steel
(Part 1: Hydrogen Absorbing Characteristic and Ductility", The Society of Materials
Science, Japan, proceedings of 51
st academic lecture meeting, 2002, vol. 8, pp17-18 and Tomohiko HOJO et. al "Influence
of Austempering Temperature on Hydrogen Embrittlement of High, for example, describe
investigations into the hydrogen embrittlement resisting property of the TRIP steel.
It is pointed out that, among the TRIP steels, TBF steel has particularly high hydrogen
absorbing capacity, and observation of a fracture surface of the TBF steel shows the
restriction of quasi cleavage fracture due to storage of hydrogen. However, the TBF
steels reported in the documents described above show delayed fracture characteristic
of about 1000 seconds at the most in terms of the time before crack occurrence measured
in cathode charging test, indicating that these steels are not meant to endure the
harsh operating environment such as that of automobile parts over a long period of
time. Moreover, since the heat treatment conditions reported in the documents described
above involve heating temperature being set higher, there are such problems as low
efficiency of practical manufacturing process. Thus it is strongly required to develop
a new species of TBF steel that provides high production efficiency as well. Also
there has been such a problem that press forming operation leads to lower hydrogen
embrittlement resisting property.
[0014] As described above, there have been virtually no TRIP steels containing residual
austenite that have been developed so as to demonstrate high workability when processed
to form parts, by taking measures to counter hydrogen embrittlement after the forming
process in consideration of the harsh operating environment such as that of automobile
parts over a long period of time.
[0015] The present invention has been made with the background described above, and has
an object of providing a high strength thin steel sheet that shows high hydrogen embrittlement
resisting property in a harsh operating environment over a long period of time after
the process of forming the steel sheet into a part, and has improved workability and
tensile strength of 1180 MPa or higher.
[0016] In order to achieve the object described above, the present inventors conducted a
research on a steel sheet that shows high hydrogen embrittlement resisting property
after the forming process, and demonstrates improved workability which is the characteristic
property of the TRIP steel sheet during the forming process. Through the research,
it was found that it is very important to control the metallurgical structure after
the forming process in order to achieve high hydrogen embrittlement resisting property
after the forming process. Specifically, it was found that it is important that the
metal structure after the stretch forming process is constituted from:
1% or more of residual austenite;
80% or more in total of bainitic ferrite and martensite; and
9% or less (may be 0%) in total of ferrite and pearlite in terms of the proportion
of area to the entire structure, wherein the mean axis ratio (major axis/minor axis)
of the residual austenite grains is 5 or higher.
A first high strength thin steel sheet having high hydrogen embrittlement resisting
property according to the present invention comprises higher than 0.25 and up to 0.60%
of C (contents of components given in terms of percentage in this patent application
all refer to percentage by weight), 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or
less P, 0.02% or less S and 1.5% or less (higher than 0%) of Al, while iron and inevitable
impurities making up the rest, wherein the metallurgical structure comprises:
1% or more residual austenite;
80% or more in total of bainitic ferrite and martensite; and
9% or less (may be 0%) in total of ferrite and pearlite in the proportion of area
to the entire structure, and wherein the mean axis ratio (major axis/minor axis) of
the residual austenite grains is 5 or higher, and the steel has tensile strength of
1180 MPa or higher.
[0017] The present inventors also conducted another research from a point of view that was
different from that of the former research, and found that high hydrogen embrittlement
resisting property after the forming process can be achieved by controlling the metal
structure after the forming process as follows. It is important that the metal structure
after the forming process comprises:
1% or more residual austenite;
the mean axis ratio (major axis/minor axis) or the residual austenite grains is 5
or higher.
mean length of minor axes of the residual austenite grains is 1 µm or less; and
minimum distance between the residual austenite grains is 1 µm or less.
When the metal structure is controlled as described above, hydrogen embrittlement
resisting property of the high strength thin steel sheet can be sufficiently improved
without adding much alloy elements. The phrase "after the forming process" means the
state of the steel sheet after being stretched with an elongation ratio of 3%. Specifically,
the steel sheet is subjected to uniaxial stretching of 3% at the room temperature
(the stretching process of 3% elongation may hereinafter be referred to simply as
"processing").
[0018] A second high strength thin steel sheet having high hydrogen embrittlement resisting
property according to the present invention comprises higher than 0.25 and up to 0.60%
of C, 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less P, 0.02% or less S, 0.5%
or less (higher than 0%) A1, while iron and inevitable impurities making up the rest,
wherein the metal structure after the stretch forming process of 3% elongation comprises:
1% or more residual austenite;
the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5
or higher;
mean length of minor axes of the residual austenite grains is 1 µm or less;
minimum distance between the residual austenite grains is 1 µm or less; and
tensile strength is 1180 MPa or higher.
According to the present invention, it is made possible to manufacture, with a high
level of productivity, a high strength thin steel sheet having tensile strength of
1180 MPa or higher that neutralizes hydrogen that infiltrates from the outside after
the steel sheet has been formed into a part thereby to maintain satisfactory hydrogen
embrittlement resisting property, and demonstrates high workability during the forming
process. Use of the high strength thin steel sheet makes it possible to manufacture
high strength parts that hardly experience delayed fracture, such as bumper, impact
beam and other reinforcement members and other automobile parts such as seat rail,
pillar, etc.
[0019]
Fig. 1 is a schematic perspective view of a part used in pressure collapse test in
Example 1.
Fig. 2 is a side view schematically showing the setup of pressure collapse test in
Example 1.
Fig. 3 is a schematic perspective view of a part used in impact resistance test in
Example 1.
Fig. 4 is a sectional view along A-A in Fig. 3.
Fig. 5 is a side view schematically showing the setup of impact resistance test in
Example 1.
Fig. 6 is a photograph of TEM observation (magnification factor 15000) of No.101 (inventive
steel) of Example 1.
Fig. 7 is a photograph of TEM observation (magnification factor 15000) of No.120 (comparative
steel) of Example 1.
Fig. 8 is a photograph of TEM observation (magnification factor 15000) of No.201 (inventive
steel) of Example 2.
Fig. 9 is a photograph of TEM observation (magnification factor 15000) of No.220 (comparative
steel) of Example 2.
Fig. 10 is a graph showing the relationship between the mean axis ratio of the residual
austenite grains and hydrogen embrittlement risk index.
Fig. 11 is a diagram schematically showing the minimum distance between residual austenite
grains.
Fig. 12 is a photograph of TEM observation (magnification factor 15000) of No.301
(inventive steel) of Example 3.
Fig. 13 is a photograph of TEM observation (magnification factor 60000) of No.301
(inventive steel) of Example 3.
Fig. 14 is a photograph of TEM observation (magnification factor 15000) of No.313
(comparative steel) of Example 3.
[0020] (First Embodiment)
The first high strength thin steel sheet according to the present invention is constituted
from higher than 0.25 and up to 0.60% of C (contents of components given in terms
of percentage in this patent application all refer to percentage by weight), 1.0 to
3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less P, 0.02% or less S, 1.5% or less (higher
than 0%) of Al, 1.0% or less (higher than 0%) of Mo and 0.1% or less (higher than
0%) of Nb, while iron and inevitable impurities making up the rest, and is characterized
in that:
(i) the metal structure after the forming process contains:
1% or more residual austenite;
80% or more in total of bainitic ferrite and martensite; and
9% or less (may be 0%) in total of ferrite and pearlite in terms of the proportion
of area to the entire structure, and
the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5
or higher; and
(ii) the steel contains a specified amount of Mo and/or Nb.
The requirements described above have reasons as follows.
[0021] (Metal structure after stretch forming by 3% elongation)
Metal structure after stretch forming process by 3% elongation was specified because,
in various experiments conducted for the actual processing conditions in the manufacture
of a part, best correlation between the result of laboratory test and the actual occurrence
of cracks in the part was observed when the part was processed by stretch forming
with an elongation ratio of 3%.
The phrase "after the forming process" means the state of the steel sheet after being
stretch formed with elongation of 3%. Specifically, the steel sheet is subjected to
elongation of 3% by uniaxial stretching at the room temperature (the stretch forming
process of 3% elongation may hereinafter be referred to simply as "process").
[0022] (1% or more residual austenite in the area proportion to the entire structure)
It is necessary that the metal contains 1% or more residual austenite in the area
proportion to the entire structure after the process of forming the part, in order
to achieve high hydrogen embrittlement resisting property in harsh operating environment
over an extended period of time after forming the part. Content of the residual austenite
is preferably 2% or higher, and more preferably 3% or higher. Since the desired level
of high strength cannot be obtained when an excessive amount of residual austenite
is contained after processing, it is recommended to set an upper limit of 20% (more
preferably 15%) to the residual austenite content.
[0023] (Mean axis ratio (major axis/minor axis) of the residual austenite grains: 5 or higher)
Lath-shaped grains of residual austenite after the process have far higher capacity
of trapping hydrogen than carbide. When the mean axis ratio (major axis/minor axis)
of the residual austenite grains is 5 or higher, in particular, it was found that
hydrogen that infiltrates from the outside through atmospheric corrosion can be substantially
neutralized thereby to achieve remarkable achievement in hydrogen embrittlement resisting
property. The mean axis ratio of the residual austenite grains is preferably 10 or
higher, and more preferably 15 or higher.
[0024] The residual austenite refers to a region that is observed as FCC (face centered
cubic lattice) by the FE-SEM/EBSP method which will be described later. Measurement
by the EBSP may be done, for example, by measuring a measurement area (about 50 by
50 µm) at an arbitrarily chosen position in a surface parallel to the rolled surface
at a position of one quarter of the thickness at measuring intervals of 0.1 µm. The
measuring surface is prepared by electrolytic polishing in order to prevent the residual
austenite from transforming. Then the test piece is set in the lens barrel of an FE-SEM
equipped with an EBSP detector (of which details will be described later) and is irradiated
with electron beam. An EBSP image projected onto a screen is captured by a high sensitivity
camera (VE-1000-SIT manufactured by Dage-MTI Inc.) and is sent to a computer. The
computer carries out image analysis and generates color mapping of the FCC phase through
comparison with a structural pattern simulated with a known crystal system (FCC (face
centered cubic lattice) phase in the case of residual austenite). Area proportion
of the region that is mapped as described above is taken as the area proportion of
the residual austenite. This analysis was carried out by means of hardware and software
of OIM (Orientation Imaging Microscopy™) system of TexSEM Laboratories Inc.
[0025] The mean axis ratio was determined by measuring the major axis and minor axis of
residual austenite crystal grains existing in each of three arbitrarily chosen fields
of view in the observation by means of TEM (transmission electron microscope) with
magnification factor of 15000, and averaging the ratios of major axis to minor axis.
[0026] (80% or more in total of bainitic ferrite and martensite)
In order to decrease the number of intergranular fracture initiating points in the
steel thereby to surely decrease the concentration of diffusive hydrogen to a harmless
level and achieve a high strength, it is desirable to form the matrix phase of the
steel structure after processing from a binary phase structure of bainitic ferrite
and martensite with the bainitic ferrite acting as the main phase, instead of the
single phase structure of martensite that is generally used for high strength steels.
[0027] In the single phase structure of martensite, a carbide (for example, film-like cementite)
is likely to precipitate in the grain boundaries, thus making intergranular fracture
likely to occur. In the case of the binary structure of bainitic ferrite and martensite
with the bainitic ferrite acting as the main phase, in contrast, the bainitic ferrite
is a hard phase and therefore it is easy to increase the strength of the entire structure
as in the case of the single phase of martensite. The hydrogen embrittlement resisting
property can also be improved as much hydrogen is trapped in the dislocations. It
also has such an advantage that coexistence of the bainitic ferrite and the residual
austenite which will be described later prevents the generation of carbide that acts
as the intergranular fracture initiating points, and it becomes easier to create the
lath-shaped residual austenite in the boundaries of lath-shaped bainitic ferrite.
[0028] Accordingly, it is required in the present invention that the binary structure of
bainitic ferrite and martensite occupy 80% or more, preferably 85% or more and more
preferably 90% or more of the entire structure after the stretch forming processing
to elongate by 3%. Upper limit of the proportion may be determined by the balance
with other structure (residual austenite), and is set to 99% when the other structures
(ferrite, etc.) than the residual austenite is not contained.
[0029] The bainitic ferrite referred to in the present invention is plate-shaped ferrite
having a lower structure of high density of dislocations. It is clearly distinguished
from polygonal ferrite that has lower structure including no or very low density of
dislocations, by SEM observation as follows.
[0030] Area proportion of bainitic ferrite structure is determined as follows. A test piece
is etched with Nital etchant. A measurement area (about 50 by 50 µm) at an arbitrarily
chosen position in a surface parallel to the rolled surface at a position of one quarter
of the thickness is observed with SEM (scanning electron microscope) (magnification
factor of 1500) thereby to determine the area proportion.
[0031] Bainitic ferrite is shown with dark gray color in SEM photograph (bainitic ferrite,
residual austenite and martensite may not be distinguishable in the case of SEM observation),
while polygonal ferrite is shown black in SEM photograph and has polygonal shape that
does not include residual austenite and martensite inside thereof.
[0032] The SEM used in the present invention is a high-resolution FE-SEM (Field Emission
type Scanning Electron Microscope XL30S-FEG manufactured by Philips Inc.) equipped
with an EBSP (Electron Back Scattering Pattern) detector, that has a merit of being
capable of analyzing the area observed by the SEM at the same time by means of the
EBSP detector. EBSP detection is carried out as follows. When the sample surface is
irradiated with electron beam, the EBSP detector analyzes the Kikuchi pattern obtained
from the reflected electrons, thereby to determine the crystal orientation at the
point where the electron beam has hit upon. Distribution of orientations over the
sample surface can be measured by scanning the electron beam two-dimensionally over
the sample surface while measuring the crystal orientation at predetermined intervals.
The EBSP detection method has such an advantage that different structures that are
regarded as the same structure in the ordinary microscopic observation but have different
crystal orientations can be distinguished by the difference in color tone.
[0033] (9% or less (may be 0%) in total of ferrite and pearlite) The steel sheet after the
processing may be constituted either from only the structures described above (namely,
a mixed structure of bainitic ferrite + martensite and residual austenite), or may
include other structure such as ferrite (the term ferrite used herein refers to polygonal
ferrite, that is a ferrite structure that includes no or very few dislocations) or
pearlite to such an extent that the effect of the present invention is not compromised.
Such additional components are structures that can inevitably remain in the manufacturing
process of the present invention, of which concentration is preferably as low as possible,
within 9%, preferably less than 5% and more preferably less than 3% according to the
present invention.
[0034] In order to maintain high hydrogen embrittlement resisting property after the forming
process, for example, large content of residual austenite of 5% or more may be contained
in the steel sheet prior to the forming process, or large amount of fine residual
austenite grains may be dispersed in the structure. Alternatively, forming process
conditions may be controlled so as to make the residual austenite less likely to transform
(for example, form the part by bending operation or control the forming temperature
and/or stretching speed). The most desirable means of improving the workability and
hydrogen embrittlement resisting property at the same time while maintaining the content
of residual austenite before and after the processing substantially constant within
an appropriate range and maintaining other properties (high strength, etc.) is to
satisfy the following requirements (A) and (B).
[0035] (A) Increase C content in the composition and increase the concentration of C in
the residual austenite.
[0036] Although residual austenite transforms into martensite when the steel sheet is deformed
(processed), high content of C in the residual austenite stabilizes it so that further
transformation becomes unlikely to occur. Thus residual austenite can be retained
after the forming process, thereby maintaining the high hydrogen embrittlement resisting
property.
[0037] According to the present invention, higher than 0.25% of C is contained in order
to achieve the effects described above. C is also an element required to achieve a
high strength of 1180 MPa or higher, and 0.27% or more, preferably 0.30% or more C
is contained. However, in order to ensure corrosion resistance, concentration of C
is limited within 0.6%, preferably 0.55% or lower and more preferably 0.50% or lower
in the present invention.
[0038] It is recommended to increase the C content in the steel sheet as described above,
thereby to maintain the concentration of C in the residual austenite (CγR) of 0.8%
or higher. Controlling the value of CyR to 0.8% or higher enables it to effectively
improve the elongation property, which is preferably 1.0% or higher and more preferably
1.2% or higher. While it is preferable that CγR is as high as possible, it is considered
that in practice there is an upper limit of around 1.6%.
[0039] (B) Form the residual austenite in fine lath-shaped grains.
[0040] Residual austenite formed in fine lath-shaped grains does not undergo excessive transformation
during the forming process, thus enabling it to maintain the residual austenite.
[0041] Some of the TRIP steels of the prior art have unsatisfactory hydrogen embrittlement
resisting property despite sufficient content of residual austenite. The reason may
be that, since residual austenite existing in the TRIP steel of the prior art generally
has block shape of size on micrometer order, it can easily transform into martensite
when being stressed and may act as the starting point of mechanical destruction. Through
a research conducted by the present inventors, it was found that residual austenite
formed in lath shape is more stable and less likely to transform into martensite than
the residual austenite of the prior art that has block shape, given the same amount
of deformation. This difference may be caused by the difference in the way in which
the stress is applied and in the difference in spatial restriction, although not fully
elucidated. Stabilization of residual austenite during processing has no influence
on the lowering of workability of TRIP steel sheet due to induced transformation.
According to the present invention, induced transformation proceeds efficiently and
high workability can be achieved without hardly reducing the residual austenite, when
the residual austenite is formed into fine lath shape as described above.
[0042] Lath-shaped grains of residual austenite having mean axis ratio (major axis/minor
axis) of 5 or higher (preferably 10 or higher, and more preferably 15 or higher) minimizes
the decrease of residual austenite during processing and makes it possible to easily
achieve mean axis ratio (major axis/minor axis) of 5 or higher after processing, put
the hydrogen absorbing capability of the residual austenite into full play and greatly
improve hydrogen embrittlement resisting property. While no upper limit of the mean
axis ratio is specified for the consideration of improvement in hydrogen embrittlement
resisting property, the residual austenite grains are required to have certain level
of thickness in order to achieve the TRIP effect during processing. Thus it is preferable
to set an upper limit to 30, more preferably to 20 or less.
[0043] According to a preferred embodiment of the present invention, Mo and Nb are added
for the purpose of reducing the size of the residual austenite grains. Mo has the
effects of strengthening the grain boundary so as to suppress hydrogen embrittlement
from occurring, in addition to reducing the size of the residual austenite grains.
Mo also has the effect of improving the hardenability of the steel sheet. It is recommended
to add 0.005% or more of Mo in order to achieve these effects. More preferably 0.1%
or more of Mo is added. However, since the effects described above reach saturation
when the Mo content exceeds 1.0%, resulting in economical disadvantage, Mo content
is limited to 0.8% or less and more preferably to 0.5% or less.
[0044] Nb, in cooperation with Mo, acts very effectively to decrease the grain size of the
structure. Nb also has the effect of increasing the strength of the steel sheet. It
is recommended to add 0.005% or more of Nb in order to achieve these effects. More
preferably 0.01% or more of Nb is added. However, since the effects described above
reach saturation when an excessive Nb content is included, resulting in economical
disadvantage, Nb content is limited to 0.1% or less and more preferably to 0.08% or
less.
[0045] In order to readily obtain the structure described above after processing, it is
recommended to make the steel sheet constituted from 80% or more (preferably 85% or
more, and more preferably 90% or more) in total of bainitic ferrite and martensite,
and 9% or less (preferably less than 5%, and more preferably less than 3% containing
0%) in total of ferrite and pearlite making up the rest of the residual austenite
before processing. This is because it is preferable that the steel sheet has high
hydrogen embrittlement resisting property prior to the processing as well as after
the processing, and this constitution makes it easier to achieve the specified strength.
[0046] While this embodiment is characterized in that metal structure is controlled after
processing, it is necessary to control the other components as described below, in
order to form the metal structure and efficiently improve hydrogen embrittlement resisting
property and strength thereby to ensure ductility required for the thin steel sheet.
[0047] <Si: 1.0 to 3.0%>
Si is an important element that effectively suppresses the residual austenite from
decomposing and carbide from being generated, and is also effective in enhancing substitution
solid solution for hardening the material. In order to make full use of these effects,
it is necessary to include Si in a concentration of 1.0% or higher, preferably 1.2%
or higher and more preferably 1.5% or higher. However, excessively high content of
Si leads to conspicuous formation of scales due to hot rolling and makes it necessary
to remove flaws, thus adding up to the manufacturing cost and resulting in economical
disadvantage. Therefore Si content is controlled within 3.0%, preferably within 2.5%
and more preferably within 2.0%.
[0048] <Mn: 1.0 to 3.5%>
Mn is an element required to stabilize austenite and obtain desired residual austenite.
In order to make full use of this effect, it is necessary to add Mn in concentration
of 1.0% or higher, preferably 1.2% or higher, and more preferably 1.5% or higher.
However, adding an excessive amount Mn leads to conspicuous segregation and poor workability.
Therefore upper limit to the concentration of Mn is set to 3.5% and more preferably
to 3.0% or less.
[0049] <P: 0.15% or lower (higher than 0%)>
P intensifies intergranular fracture due to intergranular segregation, and the content
thereof is therefore preferably as low as possible. Upper limit to the concentration
of P is set to 0.15%, preferably 0.1% or less and more preferably to 0.05% or less.
[0050] <S: 0.02% or lower (higher than 0%)>
S intensifies the absorption of hydrogen into the steel sheet in corrosive environment,
and the content thereof is therefore preferably as low as possible. Upper limit to
the concentration of S is set to 0.02%.
[0051]
<A1: 1.5% or less (higher than 0%)> (In the case of inventive steel 1)
<A1: 0.5% or less (higher than 0%)> (In the case of inventive steel 2)
0.01% or higher content of Al may be included for the purpose of deoxidation. In addition
to deoxidation, Al also has the effects of improving the corrosion resistance and
improving hydrogen embrittlement resisting property.
[0052] The mechanism of improving the corrosion resistance is supposedly based on the improvement
of corrosion resistance of the matrix phase per se and the effect of formation rust
generated by atmospheric corrosion, while the effect of the formation rust presumably
has greater contribution. This is supposedly because the formation rust is denser
and better in protective capability than ordinary iron rust, and therefore depresses
the progress of atmospheric corrosion so as to decrease the amount of hydrogen generated
by the atmospheric corrosion, thereby to effectively suppress the occurrence of hydrogen
embrittlement, and hence the delayed fracture.
[0053] While details of the mechanism of improvement of the hydrogen embrittlement resistance
by Al is not known, it is supposed that condensing of Al on the surface of the steel
makes it difficult for hydrogen to infiltrate into the steel, and the decreasing diffusion
rate of hydrogen in the steel makes it difficult for hydrogen to migrate so that hydrogen
embrittlement becomes less likely to occur. In addition, stability of lath-shaped
residual austenite improved by the addition of Al is believed to contribute to the
improvement of hydrogen embrittlement resisting property.
[0054] In order to effectively achieve the effects of Al in improving the corrosion resistance
and improving the hydrogen embrittlement resisting property, Al content is controlled
to 0.2% or higher, preferably 0.5% or higher.
[0055] However, Al content must be controlled within 1.5% in order to prevent inclusions
such as alumina from increasing in number and size so as to ensure satisfactory workability,
ensure the generation of fine residual austenite grains, suppress corrosion from proceeding
from the inclusion containing Al as the starting point, and prevent the manufacturing
cost from increasing. In view of the manufacturing process, it is preferable to control
so that A3 point is not higher than 1000°C.
[0056] As the Al content increases, inclusions such as alumina increase and workability
becomes poorer. In order to suppress the generation of the inclusions such as alumina
and make a steel sheet having higher workability, Al content is restricted within
0.5%, preferably within 0.3% and more preferably within 0.1%.
[0057] While constituent elements (C, Si, Mn, P, S, Al, Mo, Nb) of the steel of this embodiment
is as described above with the rest substantially being Fe, it may include inevitable
impurities introduced into the steel depending on the stock material, production material,
manufacturing facility and other circumstances, containing 0.001% or less of N (nitrogen).
In addition, other elements as described below may be intentionally added to such
an extent that does not adversely affect the effects of the present invention.
[0058] <B: 0.0002 to 0.01%>
B is effective in increasing the strength of the steel sheet, and it is preferable
that 0.0002% or more (more preferably 0.0005% or more) B is contained. However, an
excessive content of B leads to poor hot processing property. Therefore, it is preferable
to control the concentration of B to within 0.01% (more preferably within 0.005%).
[0059] <At least one selected from among Ca: 0.0005% to 0.005%, Mg: 0.0005% to 0.01% and
REM: 0.0005% to 0.01%)
Ca, Mg and REM (rare earth element) are effective in suppressing an increase in hydrogen
ion concentration, that is, a decrease in pH in the atmosphere of the interface due
to corrosion of the steel sheet surface, thereby to improve the corrosion resistance
of the steel sheet. These elements are also effective in controlling the form of sulfide
contained in the steel and improve the workability of the steel. In order to achieve
the effects described above, it is recommended to add each of Ca, Mg and REM in concentration
of 0.0005% or higher. However, since excessive contents of these elements leads to
poor workability, it is preferable to keep the concentration of Ca to 0.005% or less
and concentration of Mg and REM each within 0.01%.
[0060] (Second Embodiment)
The second high strength thin steel sheet according to the present invention is constituted
from higher than 0.25% and up to 0.60% of C (contents of components given in terms
of percentage in this patent application all refer to percentage by weight), 1.0 to
3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less of P, 0.02% or less of S, 1.5% or less
(higher than 0%) of Al while iron and inevitable impurities constitute the rest, wherein:
(i) the structure after the forming process comprises: 1% or more residual austenite;
mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher;
80% or more in total of bainitic ferrite and martensite; and
9% or less (may be 0%) in total of ferrite and pearlite in the proportion of area
to the entire structure, and
(ii) the steel contains specified amount of Cu and/or Ni.
[0061] The requirements (i) have the reasons as described above.
[0062] The requirement(ii) described above has the reason as follows.
[0063] Specific measures were studied to retain residual austenite after processing, control
the shape of the residual austenite grains, improve the hydrogen trapping capability
and reliably reduce the concentration of diffusive hydrogen in the steel sheet to
a harmless level by:(a) sufficiently suppressing the generation of hydrogen from the
steel sheet in corrosive environment; and (b) suppressing hydrogen that has been generated
from infiltrating the steel sheet.
[0064] It was found that it is very effective to include 0.003 to 0.5% of Cu and/or 0.003
to 1.0% of Ni in achieving the objectives of (a) and (b), and that the effect of improving
hydrogen embrittlement resisting property through control of the structure can be
achieved further by containing these elements.
[0065] Specifically, presence of Cu and Ni improves the corrosion resistance of the steel,
and effectively suppresses the generation of hydrogen due to corrosion of the steel
sheet. These elements also have the effect of promoting the generation of iron oxide,
α-FeOOH, that is believed to be particularly stable thermodynamically and have protective
property among various forms of rust generated in the atmosphere. By assisting the
generation of this rust, it is made possible to suppress hydrogen that has been generated
from infiltrating into the spring steel thereby to sufficiently improve the hydrogen
embrittlement resisting property to endure in harsh corrosive environment. This effect
can be achieved particularly satisfactorily when Cu and Ni are contained at the same
time.
[0066] In order to achieve the effects described above, concentration of Cu, if added, should
be 0.003% or higher, preferably 0.05% or higher and more preferably 0.1% or higher.
Concentration of Ni, if added, should be 0.003% or higher, preferably 0.05% or higher
and more preferably 0.1% or higher.
[0067] Since excessively high concentration of either Cu or Ni is detrimental to workability,
it is preferable to limit the Cu content to 0.5% or lower and limit the Ni content
to 1.0% or lower.
[0068] In order to achieve high hydrogen embrittlement resisting property after the forming
process by retaining the predetermined amount of residual austenite after the forming
process as in (i) described above, for example, 5% or more residual austenite may
be contained in the steel sheet prior to the forming process, or large amount of fine
residual austenite grains may be dispersed in the structure. Alternatively, forming
process conditions may be controlled so as to make the residual austenite less likely
to transform (for example, form the part by bending operation or control the forming
temperature and/or stretching speed). The most desirable means of improving the workability
and hydrogen embrittlement resisting property at the same time while maintaining the
content of residual austenite before and after the processing substantially constant
within an appropriate range and maintaining other properties (high strength, etc.)
is to satisfy the requirements (A) and (B) described previously.
[0069] While this embodiment is characterized in that metal structure is controlled after
processing and predetermined amount of Cu and/or Ni are added, it is necessary to
control the other components as described below, in order to readily form the metal
structure and efficiently improve hydrogen embrittlement resisting property and strength
thereby to ensure ductility required for the thin steel sheet.
[0070] While constituent elements (C, Si, Mn, P, S, Al, Cu and/or Ni) of the steel of this
embodiment are as described above with the rest substantially being Fe, it may include
inevitable impurities introduced into the steel depending on the stock material, production
material, manufacturing facility and other circumstances, containing 0.001% or less
of N (nitrogen). In addition, other elements as described below may be intentionally
added to such an extent that does not adversely affect the effects of the present
invention.
[0071] <Ti and/or V: 0.003 to 1.0% in total>
Ti has the effect of assisting in the generation of protective rust, similarly to
Cu and Ni. The protective rust has a very valuable effect of suppressing the generation
of β-FeOOH that appears in chloride environment and has adverse effect on the corrosion
resistance (and hence on the hydrogen embrittlement resisting property). Formation
of such a protective rust is promoted particularly by adding Ti and V (or Zr). Ti
renders the steel high corrosion resistance, and also has the effect of cleaning the
steel.
[0072] V is effective in increasing the strength of the steel sheet and decreasing the size
of crystal grains, in addition to having the effect of improving hydrogen embrittlement
resistance through cooperation with Ti, as described previously.
[0073] In order to fully achieve the effect of Ti and/or V described above, it is preferable
to add Ti and/or V in total concentration of 0.003% or higher (more preferably 0.01%
or higher). For the purpose of improving hydrogen embrittlement resisting property,
in particular, it is preferable to add more than 0.03% of Ti, more preferably 0.05%
or more of Ti. However, the effects described above reach saturation when an excessive
amount of Ti is added, resulting in economical disadvantage. Excessive V content also
increases the precipitation of much carbonitride and leads to poor workability and
lower hydrogen embrittlement resisting property. Therefore, it is preferable to control
the total concentration of Ti and/or V to within 1.0%, more preferably within 0.5%.
[0074] <Zr: 0.003 to 1.0%>
Zr is effective in increasing the strength of the steel sheet and decreasing the crystal
grain size, and also has the effect of improving hydrogen embrittlement resisting
property through cooperation with Ti. In order to sufficiently achieve these effects,
it is preferable that 0.003% or more of Zr is contained. However, excessive Zr content
increases the precipitation of carbonitride and leads to poor workability and lower
hydrogen embrittlement resisting property. Therefore, it is preferable to control
the concentration of Zr to within 1.0%.
[0075] <Mo: 1.0% or less (higher than 0%)>
Mo has the effects of stabilizing austenite so as to retain the residual austenite,
and suppress the infiltration of hydrogen thereby to improve hydrogen embrittlement
resisting property. Mo also has the effect of improving the hardenability of the steel
sheet. In addition, Mo strengthens the grain boundary so as to suppress hydrogen embrittlement
from occurring. It is recommended to add 0.005% or more Mo in order to achieve these
effects. More preferably 0.1% or more Mo is added. However, since the effects described
above reach saturation when the Mo content exceeds 1.0%, resulting in economical disadvantage,
Mo content is limited to 0.8% or less and more preferably to 0.5% or less.
[0076] <Nb: 0.1% or less (higher than 0%)>
Nb is very effective in increasing the strength of the steel sheet and decreasing
the grain size of the structure. Nb achieves these effects particularly effectively
in cooperation with Mo. In order to achieve these effects, it is recommended to include
0.005% or more of Nb. More preferably 0.01% or more of Nb is added. However, since
the effects described above reach saturation when an excessive Nb content is included,
resulting in economical disadvantage, Nb content is limited to 0.1% or less and more
preferably to 0.08% or less.
[0077] <B: 0.0002 to 0.01%>
B is effective in increasing the strength of the steel sheet, and it is preferable
that 0.0002% or more (more preferably 0.0005% or more) B is contained in order to
achieve these effects. However, an excessive content of B leads to poor hot processing
property. Therefore, it is preferable to control the concentration of B within 0.01%
(more preferably within 0.005%).
[0078] <At least one kind selected from among a group consisting of Ca: 0.0005% to 0.005%,
Mg: 0.0005% to 0.01% and REM: 0.0005% to 0.01%)
Ca, Mg and REM (rare earth element) are effective in suppressing an increase in hydrogen
ion concentration, that is, a decrease in pH in the atmosphere of the interface due
to corrosion of the steel sheet surface, thereby to improve the corrosion resistance
of the steel sheet. It is also effective in controlling the form of sulfide in the
steel and improving the workability of the steel. In order to achieve the effects
described above, it is recommended to add each of Ca, Mg and REM in concentration
of 0.0005% or higher. However, since excessive contents of these elements leads to
poor workability, it is preferable to keep the concentrations of Ca within 0.005%,
Mg and REM each within 0.01%.
[0079] (Third Embodiment)
A third high strength thin steel sheet according to the present invention is constituted
from higher than 0.25 and up to 0.60% of C (contents of components given in terms
of percentage in this patent application all refer to percentage by weight), 1.0 to
3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less of P, 0.02% or less of S, 1.5% or less
(higher than 0%) of Al, while iron and inevitable impurities making up the rest, wherein:(iii)
the structure satisfies the following requirements after forming:
1% or more residual austenite;
the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5
or higher;
mean length of minor axes of the residual austenite grains is 1 µm or less; and
minimum distance between residual austenite grains is 1 µm or less.
When the metal structure is controlled as described above, hydrogen embrittlement
resisting property of the high strength thin steel sheet can be sufficiently improved
without adding much alloy elements.
The phrase "after the forming process" means the state of the steel sheet after being
stretch formed with elongation of 3%. Specifically, the steel sheet is subjected to
elongation of 3% by uniaxial stretching at the room temperature (the stretch forming
process of 3% elongation may hereinafter be referred to simply as "process").
[0080] The requirements for the residual austenite of the present invention will now be
described in detail below.
[0081] <1% or more residual austenite>
<Mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or
higher >
It is necessary that the metal structure contains 1% or more residual austenite in
terms of area proportion to the entire structure after processing, in order to achieve
high hydrogen embrittlement resisting property in harsh operating environment over
an extended period of time after forming the part. Residual austenite contributes
not only to the improvement of hydrogen embrittlement resisting property as described
above, but also to the improvement of total elongation as has been known in the prior
art. Content of the residual austenite is preferably 2% or higher, and more preferably
3% or higher. Since the desired level of high strength cannot be obtained when an
excessive amount of residual austenite is contained, it is recommended to set an upper
limit of 15% (more preferably 10%) to the residual austenite content.
[0082] Lath-shaped grains of residual austenite after processing have far higher capacity
of trapping hydrogen than carbide. Fig. 1 is a graph showing the relationship between
the mean axis ratio of the residual austenite grains measured by a method to be described
later and hydrogen embrittlement risk index (measured by a method to be described
later in an example, lower value of this index means better hydrogen embrittlement
resisting property). From Fig. 1, it can be seen that hydrogen embrittlement risk
index sharply decreases when the mean axis ratio (major axis/minor axis) of the residual
austenite grains increases beyond 5. This is supposedly because, when the mean axis
ratio of the residual austenite grains becomes 5 or higher, intrinsic capability of
the residual austenite to absorb hydrogen is put into full play, so that the residual
austenite attains far higher capacity of trapping hydrogen than carbide and substantially
neutralizes the hydrogen that infiltrates from the outside through atmospheric corrosion
thereby to achieve remarkable achievement in hydrogen embrittlement resisting property.
The mean axis ratio of the residual austenite grains is preferably 10 or higher, and
more preferably 15 or higher.
[0083] <Niean length of minor axes of the residual austenite grains is 1 µm or less>
According to the present invention, it has been found that hydrogen embrittlement
resisting property can be effectively improved by dispersing fine grains of residual
austenite of lath shape. Specifically, hydrogen embrittlement resisting property can
be surely improved by dispersing the lath-shape grains of residual austenite having
sizes of 1 µm or less (submicrometer order). This is supposedly because surface area
of the residual austenite grains (interface) increases resulting in larger hydrogen
trapping capability, when larger number of fine lath-shape grains of residual austenite
having smaller mean length of minor axis are dispersed. Mean length of minor axes
of the residual austenite grains is preferably 0.5 µm or less, more preferably 0.25
µm or less.
[0084] According to the present invention, hydrogen trapping capability of the fine lath-shape
grains of residual austenite can be made far greater than that in the case of dispersing
carbide, and thereby to substantially neutralize hydrogen that infiltrates from the
outside through atmospheric corrosion, even when the same proportion by volume of
residual austenite is contained, by controlling the mean axis ratio and mean length
of minor axes of the residual austenite grains as described above.
[0085] <Minimum distance between residual austenite grains is 1 µm or less>
According to the present invention, it has been found that hydrogen embrittlement
resisting property can be improved further by controlling the minimum distance between
adjacent residual austenite grains, in addition to the above. Specifically, hydrogen
embrittlement resistance can be surely improved when the minimum distance between
residual austenite grains is 1 µm or less. This is supposedly because propagation
of cracks is suppressed so that the structure demonstrates higher resistance against
fracture, when a large number of fine lath-shape grains of residual austenite are
dispersed in proximity to each other. Minimum distance between adjacent residual austenite
grains is preferably 0.8 µm or less, and more preferably 0.5 µm or less.
[0086] The residual austenite refers to a region that is observed as FCC (face centered
cubic lattice) by the FE-SEM/EBSP method which will be described later. Measurement
by the EBSP may be done, for example, by measuring a measurement area (about 50 by
50 µm) at an arbitrarily chosen position in a surface parallel to the rolled surface
at a position of one quarter of the thickness at measuring intervals of 0.1 µm. The
measuring surface is prepared by electrolytic polishing in order to prevent the residual
austenite from transforming. Then the test piece is set in the lens barrel of an FE-SEM
equipped with the EBSP detector (of which details will be described later) and is
irradiated with electron beam. An EBSP image projected onto a screen is captured by
a high sensitivity camera (VE-1000-SIT manufactured by Dage-MTI Inc.) and is sent
to a computer. The computer carries out image analysis and generates color mapping
of the FCC phase through comparison with a structural pattern simulated with a known
crystal system (FCC (face centered cubic lattice) phase in the case of residual austenite).
Area proportion of the region that is mapped as described above is taken as the area
proportion of the residual austenite. This analysis was carried out by means of hardware
and software of OIM (Orientation Imaging Microscopy™) system of TexSEM Laboratories
Inc.
[0087] The mean axis ratio, mean length of minor axes and minimum distance between residual
austenite grains were determined as follows. The mean axis ratio of the residual austenite
grains was determined by measuring the major axis and minor axis of residual austenite
crystal grain existing in each of three arbitrarily chosen fields of view in the observation
by means of TEM (transmission electron microscope) with magnification factor of 15000,
and averaging the ratios of major axis to minor axis. The mean length of minor axes
of the residual austenite grains was determined by averaging the lengths of minor
axes measured as described above. The minimum distance between adjacent residual austenite
grains was determined by measuring the distance between adjacent residual austenite
grains that were aligned in the direction of major axis as shown as (a) in Fig. 2
(distance (b) in Fig. 2 is not regarded as the minimum distance between the grains)
by observing with TEM (magnification factor of 15000) in each of three arbitrarily
chosen fields of view and averaging the distances measured in the three fields of
view.
[0088] In order to decrease the number of intergranular fracture initiating points in the
steel thereby to surely decrease the concentration of diffusive hydrogen to a harmless
level and achieve a high strength, it is desirable to form the matrix phase of the
steel sheet after processing from a binary phase structure of bainitic ferrite and
martensite with the bainitic ferrite acting as the main phase, instead of the single
phase structure of martensite that is generally used for high strength steels.
[0089] In the case of single phase structure of martensite, a carbide (for example, film-like
cementite) is likely to precipitate in the grain boundaries, thus making intergranular
fracture likely to occur. In the case of the binary phase structure of bainitic ferrite
and martensite with the bainitic ferrite acting as the main phase, in contrast, the
bainitic ferrite is a hard phase and therefore it is easy to increase the strength
of the entire structure as in the case of the single phase of martensite. The hydrogen
embrittlement resisting property can also be improved as much hydrogen is trapped
in the dislocations. It also has such an advantage that coexistence of the bainitic
ferrite and residual austenite which will be described later prevents the generation
of carbide that acts as the intergranular fracture initiating points, and it becomes
easier to create the lath-shaped residual austenite in the boundaries of lath-shaped
bainitic ferrite.
[0090] Accordingly, it is required in the present invention that the binary phase structure
of bainitic ferrite and martensite occupy 80% or more, preferably 85% or more and
more preferably 90% or more of the entire structure after the stretch forming processing
to elongate by 3%. Upper limit of the proportion may be determined by the balance
with other structure (residual austenite), and is set to 99% when other structure
(ferrite, etc.) than the residual austenite is not contained.
[0091] The bainitic ferrite referred to in the present invention is plate-shaped ferrite
having a lower structure of high density of dislocations. It is clearly distinguished
from polygonal ferrite that has lower structure including no or very low density of
dislocations, by SEM observation as follows.
[0092] Area proportion of bainitic ferrite structure is determined as follows. A test piece
is etched with Nital etchant. A measurement area (about 50 by 50 µm) at an arbitrarily
chosen position in a surface parallel to the rolled surface at a position of one quarter
of the thickness is observed with SEM (scanning electron microscope) (magnification
factor of 1500) thereby to determine the area proportion.
[0093] Bainitic ferrite is shown with dark gray color in SEM photograph (bainitic ferrite,
residual austenite and martensite may not be distinguishable in the case of SEM observation),
while polygonal ferrite is shown black in SEM photograph and has polygonal shape that
does not contain residual austenite and martensite inside thereof.
[0094] The SEM used in the present invention is a high-resolution FE-SEM (Field Emission
type Scanning Electron Microscope XL30S-FEG manufactured by Philips Inc.) equipped
with an EBSP (Electron Back Scatter diffraction Pattern) detector, that has a merit
of being capable of analyzing the area observed by the SEM at the same time by means
of the EBSP detector. EBSP detection is carried out as follows. When the sample surface
is irradiated with electron beam, the EBSP detector analyzes the Kikuchi pattern obtained
from the reflected electrons, thereby to determine the crystal orientation at the
point where the electron beam has hit upon. Distribution of orientations over the
sample surface can be measured by scanning the electron beam two-dimensionally over
the sample surface while measuring the crystal orientation at predetermined intervals.
The EBSP detection method has such an advantage that different structures that are
regarded as the same structure in the ordinary microscopic observation but have different
crystal orientations can be distinguished by the difference in color tone.
[0095] The metal structure after the processing may be constituted either from only the
structures described above (namely, a mixed structure of bainitic ferrite + martensite
and residual austenite), or may include other structure such as ferrite (the term
ferrite used herein refers to polygonal ferrite, that is a ferrite structure that
includes no or very few dislocations) or pearlite to such an extent that the effect
of the present invention is not compromised. Such additional components are structures
that can inevitably remain in the manufacturing process of the present invention,
of which concentration is preferably as low as possible, within 9%, preferably less
than 5% and more preferably less than 3% according to the present invention.
[0096] In order to maintain high hydrogen embrittlement resisting property after the forming
process, for example, high proportion of residual austenite, 5% or more, may be contained
in the steel sheet prior to the forming process, or a large amount of fine residual
austenite grains may be dispersed in the structure. Alternatively, forming process
conditions may be controlled so as to make the residual austenite less likely to transform
(for example, form the part by bending operation or control the forming temperature
and/or stretching speed). The most desirable means of improving the workability and
hydrogen embrittlement resisting property at the same time while maintaining the content
of residual austenite before and after the processing substantially constant within
an appropriate range and maintaining other properties (high strength, etc.) is to
satisfy the requirements (A) and (B) described previously.
[0097] While this embodiment is characterized in that the metal structure is controlled
after processing, it is necessary to control the other components as described previously,
in order to form the metal structure and efficiently improve hydrogen embrittlement
resisting property and strength thereby to ensure the level of ductility required
for the thin steel sheet.
[0098] While the present invention does not specify the manufacturing conditions, it is
recommended to apply heat treatment in the following procedure after hot rolling or
cold rolling conducted thereafter, in order to form the structure described above
that can be easily worked and has high strength and high hydrogen embrittlement resistance
after the processing, by using the steel material of the composition described above.
The recommended procedure is to keep the steel the composition described above at
a temperature (T1) in a range from A3 point to (A3 point + 50°C) for a period of 10
to 1800 seconds (t1), cool down the steel at a mean cooling rate of 3°C/s or higher
to a temperature (T2) in a range from Ms point to Bs point and keep the material at
this temperature for a period of 60 to 3600 seconds (t2).
[0099] It is not desirable that the temperature T1 becomes higher than (A3 point + 50°C)
or the period t1 is longer than 1800 seconds, in which case austenite grains grow
resulting in poor workability (elongation flanging property). When the temperature
T1 is lower than A3 point, on the other hand, desirable bainitic ferrite structure
cannot be obtained. When the period t1 is shorter than 10 seconds, austenitization
does not proceed sufficiently and therefore cementite and other alloy carbides remain.
The period t1 is preferably in a range from 30 to 600 seconds, more preferably from
60 to 400 seconds.
[0100] Then the steel sheet is cooled down. The steel is cooled at a mean cooling rate of
3°C/s or higher, for the purpose of preventing pearlite structure from being generated
while avoiding the pearlite transformation region. The mean cooling rate should be
as high as possible, and is preferably 5°C/s or higher, and more preferably 10°C/s
or higher.
[0101] After quenching to the temperature between Ms point and Bs point at the rate described
above, the steel is subjected to isothermal transformation so as to transform the
matrix phase into binary phase structure of bainitic ferrite and martensite. When
the heat retaining temperature T2 is higher than Bs, much pearlite that is not desirable
for the present invention is formed, thus hampering the formation of the predetermined
bainitic ferrite structure. When T2 is below Ms, on the other hand, the amount of
residual austenite decreases.
[0102] When the temperature holding period t2 is longer than 1800 seconds, density of dislocations
in bainitic ferrite becomes low, the amount of trapped hydrogen decreases and the
desired residual austenite cannot be obtained. When t2 is less than 60 second, on
the other hand, desired bainitic ferrite structure cannot be obtained. The length
of t2 is preferably from 90 to 1200 seconds, and more preferably from 120 to 600 seconds.
There is no restriction on the method of cooling after maintaining the heating temperature,
and air cooling, quenching or air-assisted water cooling may be employed.
[0103] In the practical manufacturing process, the annealing process described above can
be carried out easily by employing a continuous annealing facility or a batch annealing
facility. In case that a cold rolled sheet is plated with zinc by hot dipping, the
heat treatment process may be replaced by the plating process by setting the plating
conditions so as to satisfy the heat treatment conditions. The plating may also be
alloyed.
[0104] There is no restriction on the hot rolling process (or cold rolling process as required)
that precedes the continuous annealing process described above, and commonly employed
process conditions may be used. Specifically, the hot rolling process may be carried
out in such a procedure as, after hot rolling at a temperature above Ar3 point, the
steel sheet is cooled at a mean cooling rate of about 30°C/s and is wound up at a
temperature from about 500 to 600°C. In case that the hot rolled steel sheet has unsatisfactory
appearance, cold rolling may be applied in order to rectify the appearance. It is
recommended to set the cold rolling ratio in a range from 1 to 70%. Cold rolling beyond
70% leads to excessive rolling load that makes it difficult to carry out the cold
rolling.
[0105] While the present invention is addressed to thin steel sheet, there is no limitation
to the form of product, and may be applied, in addition to steel sheet made by hot
rolling or steel sheet made by cold rolling, to those subjected to annealing after
hot rolling or cold rolling, followed by chemical conversion treatment, hot-dip coating,
electroplating, vapor deposition, painting, priming for painting, organic coating
treatment or the like.
[0106] The plating process may be either galvanizing or aluminum plating. The method of
plating may be either hot-dip coating or electroplating, and the plating process may
also be followed by alloying heat treatment or multi-layer plating. A steel sheet,
that is plated or not plated, may also be laminated with a film.
[0107] . When the coating operation described above is carried out, chemical conversion
treatment such as phosphating or electrodepositing coating may be applied in accordance
to the application. The coating material may be a known resin that can be used in
combination with a known hardening agent such as epoxy resin, fluorocarbon resin,
silicone acrylic resin, polyurethane resin, acrylic resin, polyester resin, phenol
resin, alkyd resin, or melamine resin. Among these, epoxy resin, fluorocarbon resin
or silicone acrylic resin is preferably used in consideration of corrosion resistance.
Known additives that are added to coating materials such as coloring agent, coupling
agent, leveling agent, sensitization agent, antioxidant agent, anti-UV protection
agent, flame retarding agent or the like may be used.
[0108] There is also no restriction on the coating and solvent-based coating, powder coating,
water-based coating, water-dispersed coating, electrodeposition coating or like may
be employed. Desired coating layer of the coating material described above can be
formed on the steel by a known technique such as dipping, roll coater, spraying, or
curtain flow coater. The coating layer may have any proper thickness.
[0109] The high strength thin steel sheet of the present invention may be applied to high-strength
automotive components such as bumper, door impact beam, pillar and other reinforcement
members and interior parts such as seat rail, etc. Automobile components that are
manufactured by forming process also have sufficient properties (strength) and high
hydrogen embrittlement resisting property.
[0110] The present invention will now be described below by way of examples, but the present
invention is not limited to the examples. Various modifications may be conceived without
departing from the technical scope of the present invention.
[Example 1]
[0111] Sample steels A-1 through Y-1 having the compositions described in Table 1 were melt-refined
in vacuum to make test slabs. The slabs were processed in the following procedure
(hot rolling → cold rolling →continuous annealing) thereby to obtain hot-rolled steel
plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales
from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.
[0112] <Hot rolling> Starting temperature (SRT): Held at a temperature between 1150 and
1250°C for 30 minutes. Finishing temperature (FDT): 850°C
Cooling rate: 40°C/s
Winding-up temperature: 550°C
<Cold rolling> Rolling ratio: 50%
<Continuous annealing> Each steel specimen was kept at a temperature of A3 point +
30°C for 120 seconds, then cooled in air at a mean cooling rate of 20°C/s to temperature
T0 shown in Table 2, and was kept at T0 for 240 seconds, followed by air-assisted
water cooling to the room temperature.
[0113] No. 116 shown in Table 2 was made by heating a cold-rolled steel sheet to 830°C,
keeping at this temperature for 5 minutes followed by quenching in water and tempering
at 300°C for 10 minutes, thereby to form a martensite steel as a comparative example
of the high-strength steel of the prior art. No. 120 was made by heating a cold-rolled
steel sheet to 800°C, keeping at this temperature for 120 seconds, cooling down at
a mean cooling rate of 20°C/s to 350°C and keeping at this temperature for 240 seconds.
[0114] JIS No. 5 test pieces were prepared from the steel sheets obtained as described above,
and were subjected to stretch forming process with elongation of 3% mimicking the
actual manufacturing process. Metal structures of the test pieces were observed before
and after the processing, tensile strength (TS) and elongation (total elongation E1)
before the processing and hydrogen embrittlement resisting property after the processing
were measured by the following procedures.
[0115] Observation of metal structure
Metal structures of the test pieces were observed before and after the processing
as follows. A measurement area (about 50 by 50 µm) at an arbitrarily chosen position
in a surface parallel to the rolled surface at a position of one quarter of the thickness
was photographed at measuring intervals of 0.1 µm, and area proportions of bainitic
ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by
the method described previously. Then similar measurements were made in two fields
of view that were arbitrarily selected, and the measured values were averaged. Area
proportions of other structures (ferrite, pearlite, etc.) were subtracted from the
entire structure.
[0116] Mean axis ratio of the residual austenite grains of the steel sheet before and after
the processing were measured by the method described previously. Test pieces having
mean axis ratio of 5 or higher were regarded to satisfy the requirements of the present
invention (o), and those having mean axis ratio of lower than 5 were regarded to fail
to satisfy the requirements of the present invention (x).
[0117] Measurement of tensile strength (TS) and elongation (E1)
Tensile test was conducted on the JIS No. 5 test piece before processing, so as to
measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile
test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa
as measured by the method described previously, those which showed elongation of 10%
or more were evaluated as high in elongation property.
[0118] Evaluation of hydrogen embrittlement resisting property
In order to evaluate hydrogen embrittlement resisting property, the JIS No. 5 test
piece was stretched so as to elongate by 3%. Then after bending with a radius of curvature
of 15 mm, load of 1000 MPa was applied and the test piece was immersed in 5% solution
of hydrochloric acid, and the time before crack occurred was measured.
[0119] Hydrogen-charged 4-point bending test was also conducted for some steel species.
Specifically, a rectangular test piece measuring 65 mm by 10 mm made of each steel
sheet elongated by 3% was immersed in a solution of 0.5 mol of H
2SO
4 and 0.01 mol of KSCN and was subjected to cathode hydrogen charging. Maximum stress
endured without breaking for 3 hours was determined as the critical fracture stress
(DFL).
[0120] Results of these tests are shown in Table 2.
[0121]

[0122]

[0123] The results shown in Tables 1 and 2 can be interpreted as follows (numbers in the
following description are test Nos. in Table 2).
[0124] Test pieces Nos. 101 through 113 (inventive steel sheets 2) and test pieces Nos.
121 through 125 (inventive steel sheets 1) that satisfy the requirements of the present
invention have high strength of 1180 MPa or higher, and high hydrogen embrittlement
resisting property in harsh environment after the forming process. They also have
high elongation property required of the TRIP steel sheet, thus providing steel sheets
best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere.
Test pieces Nos. 121 through 125, in particular, show even better hydrogen embrittlement
resisting property.
[0125] Test pieces Nos. 114 through 120 and 126 that do not satisfy the requirements of
the present invention, in contrast, have the following drawbacks.
[0126] No. 114 made of steel species N-1 that includes insufficient C content does not have
good workability.
[0127] No. 115 made of steel species O-1 that includes insufficient Mn content does not
retain sufficient residual austenite and is inferior in hydrogen embrittlement resisting
property after the processing.
[0128] No. 116, martensite steel that is a conventional high strength steel made of steel
species P-1 that includes insufficient Si content, hardly contains residual austenite
and is inferior in hydrogen embrittlement resisting property.
It also does not show the elongation property required of a thin steel sheet.
[0129] No. 117 made of steel species Q-1 that includes excessive C content has precipitation
of carbide and is inferior in both forming workability and hydrogen embrittlement
resisting property after processing.
[0130] No. 118 made of steel species R-1 that includes excessive Mo content and No. 119
made of steel species S-1 that includes excessive Nb content are inferior in forming
workability. Nos. 118 and 119 could not undergo the processing, making it impossible
to investigate the property after the processing.
[0131] No. 120, that was made of a steel that has the composition specified in the present
invention but was not manufactured under the recommended conditions, resulted in the
conventional TRIP steel. As a result, the residual austenite does not have the mean
axis ratio specified in the present invention, while the matrix phase is not formed
in binary phase structure of bainitic ferrite and martensite, and therefore sufficient
level of hydrogen embrittlement resisting property is not achieved.
[0132] No. 126 includes Al content higher than that specified for the inventive steel sheet
1. As a result, although the predetermined amount of residual austenite is retained,
the residual austenite does not have the mean axis ratio specified in the present
invention, the desired matrix phase is not obtained and inclusions such as AlN are
generated thus resulting in poor hydrogen embrittlement resisting property.
[0133] Then parts were made by using steel species A-1, J-1 shown in Table 1 and comparative
steel sheet (590 MPa class high strength steel sheet of the prior art). Performance
(pressure collapse resistance and impact resistance) of the formed test piece were
studied by conducting pressure collapse test and impact resistance test as follows.
[0134] Pressure collapse test
The part 1 (hat channel as test piece) shown in Fig. 1 was made by using steel species
A-1, J-1 shown in Table 1 and the comparative steel sheet, and was subjected to pressure
collapse test. The part was spot welded at the positions 2 of the part shown in Fig.
1 at 35 mm intervals as shown in Fig. 1 by supplying electric current of a magnitude
less than the expulsion generating current by 0.5 kA from an electrode measuring 6
mm in diameter at the distal end. Then a die 3 was pressed against the part 1 from
above the mid portion thereof in the longitudinal direction as shown in Fig. 2, and
the maximum tolerable load was determined. Absorbed energy was determined from the
area under the load-deformation curve. The results are shown in Table 3.
[0135]
Table 3
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Maximum load |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kN) |
(kJ) |
Symbol A-1 |
1510 |
15 |
9 |
14.1 |
0.72 |
Symbol J-1 |
1491 |
12 |
11 |
14 |
0.68 |
Comparative steel sheet |
613 |
22 |
0 |
5.7 |
0.33 |
[0136] From Table 3, it can be seen that the part (test piece) made from the steel sheet
of the present invention has higher load bearing capability and absorbs greater energy
than a part made of the conventional steel sheet having lower strength, thus showing
high pressure collapse resistance.
[0137] Impact resistance test
The parts 4 (hat channel as test piece) shown in Fig. 3 were made by using steel species
A-1, J-1 shown in Table 1 and the comparative steel sheet, and were subjected to impact
resistance test. Fig. 4 is a sectional view along A-A of the part 4 shown in Fig.
3. In the impact resistance test, after the part was spot welded at the positions
5 of the part 4 similarly to the pressure collapse test, the part 4 was placed on
a base 7 as schematically shown in Fig. 5. A weight 6 (weighing 10kg) was dropped
onto the part 4 from a height of 11 meters, and the energy absorbed before the part
4 underwent deformation of 40 mm in the direction of height. The results are shown
in Table 4.
[0138]
Table 4
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kJ) |
Symbol A-1 |
1510 |
15 |
9 |
6.94 |
Symbol J-1 |
1491 |
12 |
11 |
6.65 |
Comparative steel sheet |
613 |
22 |
0 |
3.56 |
[0139] From Table 4, it can be seen that the part (test piece) made from the steel sheet
of the present invention absorbs greater energy than a part made of the conventional
steel sheet that has lower strength, thus showing higher impact resistance.
[0140] TEM photograph of the test piece made in this example is shown as reference. Fig.
6 is a photograph of TEM observation of No. 101 of the present invention. From Fig.
6, it can be seen that the high strength thin steel sheet of the present invention
contains lath-shaped residual austenite (black portion of bar shape in Fig. 6) specified
in the present invention dispersed therein. Fig. 7 is a photograph of TEM observation
of No. 120 of a comparative example. From Fig. 7, it can be seen that the high strength
thin steel sheet of No. 120 contains residual austenite (black portion of somewhat
round shape in Fig. 7), although the residual austenite has a block shape that does
not satisfy the requirements of the present invention.
[Example 2]
[0141] Sample steels A-2 through Y-2 having the compositions described in Table 5 were melt-refined
in vacuum to make test slabs. The slabs were processed in the following procedure
(hot rolling→cold rolling→continuous annealing) thereby to obtain hot-rolled steel
plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales
from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.
[0142] <Hot rolling>
Starting temperature (SRT): Held at a temperature between 1150 and 1250°C for 30 minutes.
Finishing temperature (FDT): 850°C
Cooling rate: 40°C/s
Winding-up temperature: 550°C
<Cold rolling>
Rolling ratio: 50%
<Continuous annealing>
Each steel specimen was kept at a temperature of A3 point + 30°C for 120 seconds,
then rapidly cooled (air cooling) at a mean cooling rate of 20°C/s to temperature
T0 shown in Table 2, and was kept at T0 for 240 seconds, followed by air-assisted
water cooling to the room temperature.
[0143] No. 217 in Table 6 was made by heating a cold-rolled steel sheet to 830°C, keeping
at this temperature for 5 minutes followed by quenching in water and tempering at
300°C for 10 minutes, thereby to form a martensite steel as a comparative example
of the high-strength steel of the prior art. No. 220 was made by heating a cold-rolled
steel sheet to 800°C, keeping at this temperature for 120 seconds, cooling down at
a mean cooling rate of 20°C/s to 350°C and keeping at this temperature for 240 seconds.
[0144] JIS No. 5 test pieces were prepared from the steel sheets obtained as described above,
and were subjected to stretch forming process with elongation of 3% mimicking the
actual manufacturing process. Metal structures of the test pieces were observed before
and after the processing, tensile strength (TS) and elongation (total elongation E1)
before the processing and hydrogen embrittlement resisting property after the processing
were measured by the following procedures.
[0145] Observation of metal structure
Metal structures of the test pieces were observed before and after the processing
as follows. A measurement area (about 50 by 50 µm) at an arbitrarily chosen position
in a surface parallel to the rolled surface at a position of one quarter of the thickness
was photographed at measuring intervals of 0.1 µm, and area proportions of bainitic
ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by
the method described previously. Then similar measurements were made in two fields
of view that were arbitrarily selected, and the measured values were averaged. Area
proportions of other structures (ferrite, pearlite, etc.) were subtracted from the
entire structure.
[0146] Mean axis ratio of the residual austenite grains of the steel sheet before and after
the processing were measured by the method described previously. Test pieces having
mean axis ratio of 5 or higher were regarded to satisfy the requirements of the present
invention (o), and those having mean axis ratio of lower than 5 were regarded to fail
to satisfy the requirements of the present invention (×).
[0147] Measurement of tensile strength (TS) and elongation (E1)
Tensile test was conducted on the JIS No. 5 test piece before processing, so as to
measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile
test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa
as measured by the method described previously, those which showed elongation of 10%
or more were evaluated as high in elongation property.
[0148] Evaluation of hydrogen embrittlement resisting property
In order to evaluate the hydrogen embrittlement resisting property, the JIS No. 5
test piece was stretched so as to elongate by 3%. Then after bending with a radius
of curvature of 15 mm, load of 1000 MPa was applied and the test piece was immersed
in 5% solution of hydrochloric acid, to measure the time before crack occurred.
[0149] The bent test pieces prepared as described above were subjected to accelerated exposure
test in which 3% solution of NaCl was sprayed once every day for 30 days simulating
the actual operating environment, and the number of days before crack occurred was
determined.
[0150] Hydrogen-charged 4-point bending test was also conducted for some steel species.
Specifically, a rectangular test piece measuring 65 mm by 10 mm made of each steel
sheet elongated by 3% was immersed in a solution of 0.5 mol of H
2SO
4 and 0.01 mol of KSCN and was subjected to cathode hydrogen charging.
Maximum stress endured without breaking for 3 hours was determined as the critical
fracture stress (DFL). Then the ratio (DFL ratio) of this value to the value of DFL
of test No. 203 (steel species C-2) shown in Table 6 was determined.
Results of these tests are shown in Table 6.
[0151]

[0152]

[0153] The results shown in Tables 5 and 6 can be interpreted as follows (numbers in the
following description are test Nos. in Table 6).
[0154] Test pieces Nos. 201 through 214 (inventive steel sheets 2) and test pieces Nos.
221 through 225 (inventive steel sheets 1) that satisfy the requirements of the present
invention have high strength of 1180 MPa or higher, and high hydrogen embrittlement
resisting property in harsh environment after the forming process. They also have
high elongation property required of the TRIP steel sheet, thus providing steel sheets
best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere.
Test pieces Nos. 221 through 225, in particular, show even better hydrogen embrittlement
resisting property.
[0155] Test pieces Nos. 215 through 220 and 226 that do not satisfy the requirements of
the present invention, in contrast, have the following drawbacks.
[0156] No. 215 made of steel species O-2 that includes insufficient C content has the amount
of residual austenite significantly decreased after the processing, and fails to show
the required level of hydrogen embrittlement resisting property of the present invention.
[0157] No. 216 made of steel species P-2 that includes insufficient Mn content does not
retain sufficient residual austenite and is inferior in hydrogen embrittlement resisting
property after the processing.
[0158] No. 217, martensite steel that is a conventional high strength steel made of steel
species Q-2 that includes insufficient Si content, hardly contains residual austenite
and is inferior in hydrogen embrittlement resisting property. It also does not show
the elongation property required of a thin steel sheet.
[0159] No. 218 made of steel species R-2 that includes excessive C content has precipitation
of carbide and is inferior in both the forming workability and the hydrogen embrittlement
resisting property after processing.
[0160] No. 219 made of steel species S-2 that does not include Cu and/or Ni shows insufficient
corrosion resistance and fails to show the required level of hydrogen embrittlement
resisting property of the present invention.
[0161] No. 220, that was made of a steel that has the composition specified in the present
invention but was not manufactured under the recommended conditions, resulted in the
conventional TRIP steel. As a result, the residual austenite does not have the mean
axis ratio specified in the present invention, while the matrix phase is not formed
in binary phase structure of bainitic ferrite and martensite, and therefore sufficient
level of hydrogen embrittlement resisting property is not achieved.
[0162] No. 226 includes A1 content higher than that specified for the inventive steel sheet
1. As a result, although the predetermined amount of residual austenite is retained,
the residual austenite does not have the mean axis ratio specified in the present
invention, the desired matrix phase is not obtained and inclusions such as AlN are
generated thus resulting in poor hydrogen embrittlement resisting property.
[0163] Then parts were made by using steel species A-2, K-2 shown in Table 5 and comparative
steel sheet (590 MPa class high strength steel sheet of the prior art). Performance
(pressure collapse resistance and impact resistance) of the formed test piece were
studied by conducting pressure collapse test and impact resistance test as follows.
[0164] Pressure collapse test
Maximum tolerable load was determined similarly to Example 1 by using steel species
A-2, K-2 shown in Table 5 and the comparative steel sheet. Absorbed energy was determined
from the area lying under the load-deformation curve. The results are shown in Table
7.
[0165]
Table 7
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Maximum load |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kN) |
(kJ) |
Symbol A-2 |
1512 |
14 |
13 |
14.1 |
0.7 |
Symbol K-2 |
1485 |
14 |
13 |
13.9 |
0.68 |
Comparative steel sheet |
613 |
22 |
0 |
5.7 |
0.33 |
[0166] From Table 7, it can be seen that the part (test piece) made from the steel sheet
of the present invention has higher load bearing capability and absorbs greater energy
than a part made of the conventional steel sheet that has lower strength, thus showing
higher pressure collapse resistance.
[0167] Impact resistance test
The impact resistance test was conducted similarly to Example 1 on the steel sheets
made of steel species A-2, K-2 shown in Table 5 and the comparative steel sheet. The
results are shown in Table 8.
[0168]
[Table 8]
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kJ) |
Symbol A-2 |
1512 |
14 |
13 |
7.06 |
Symbol K-2 |
1485 |
14 |
13 |
6.92 |
Comparative steel sheet |
613 |
22 |
0 |
3.56 |
[0169] From Table 8, it can be seen that the part (test piece) made from the steel sheet
of the present invention absorbs greater energy than a part made of the conventional
steel sheet having lower strength, thus showing higher impact resistance.
[0170] TEM photograph of the test piece made in this example is shown as reference. Fig.
8 is a photograph of TEM observation of No. 201 of the present invention. From Fig.
8, it can be seen that the high strength thin steel sheet of the present invention
contains lath-shaped residual austenite (black portion of bar shape in Fig. 8) specified
in the present invention dispersed therein. Fig. 9 is a photograph of TEM observation
of No. 220 of a comparative example. From Fig. 9, it can be seen that the high strength
thin steel sheet of No. 220 contains residual austenite (black portion of somewhat
round shape in Fig. 9), although the residual austenite has a block shape that does
not satisfy the requirements of the present invention.
[Example 3]
[0171] Sample steels A-3 through Q-3 having the compositions shown in Table 9 were melt-refined
in vacuum to make test slabs. The slabs were processed in the following procedure
(hot rolling→cold rolling→continuous annealing) thereby to obtain hot-rolled steel
plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales
from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.
[0172] <Hot rolling> Starting temperature (SRT): Held at a temperature between 1150 and
1250°C for 30 minutes.
Finishing temperature (FDT): 850°C
Cooling rate: 40°C/s
Winding-up temperature: 550°C
<Cold rolling> Rolling ratio: 50%
<Continuous annealing> Each steel specimen was kept at a temperature of A3 point +
30°C for 120 seconds, then rapidly cooled (air cooling) at a mean cooling rate of
20°C/s to temperature T0 shown in Table 2, and was kept at T0 for 240 seconds, followed
by air-assisted water cooling to the room temperature.
[0173] No. 311 shown in Table 10 was made by heating a cold-rolled steel sheet to 830°C,
keeping at this temperature for 5 minutes followed by quenching in water and tempering
at 300°C for 10 minutes, thereby to form a martensite steel as a comparative example
of the high-strength steel of the prior art. No. 312 was made by heating a cold-rolled
steel sheet to 800°C, keeping at this temperature for 120 seconds, cooling at a mean
cooling rate of 20°C/s down to 350°C and keeping at this temperature for 240 seconds.
[0174] JIS No. 5 test pieces were prepared from the steel sheets obtained as described above,
and were subjected to stretch forming process with elongation of 3% mimicking the
actual manufacturing process. Metal structures of the test pieces were observed before
and after the processing, tensile strength (TS) and elongation (total elongation E1)
before the processing and hydrogen embrittlement resisting property after the processing
were measured by the following procedures.
[0175] Observation of metal structure
Metal structures of the test pieces were observed before and after the processing
as follows. A measurement area (about 50 by 50 µm) at an arbitrarily chosen position
in a surface parallel to the rolled surface at a position of one quarter of the thickness
was photographed at measuring intervals of 0.1 µm, and area proportions of bainitic
ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by
the method described previously. Then similar measurements were made in two fields
of view that were arbitrarily selected, and the measured values were averaged. Area
proportions of other structures (ferrite, pearlite, etc.) were subtracted from the
entire structure.
[0176] Mean axis ratio, mean length of minor axes and minimum distance between the residual
austenite grains of the steel sheet before and after the processing were measured
by the method described previously. Test pieces having mean axis ratio of 5 or higher
were regarded to satisfy the requirements of the present invention (o), and those
having mean axis ratio of lower than 5 were regarded to fail to satisfy the requirements
of the present invention (x).
[0177] Measurement of tensile strength (TS) and elongation (E1)
Tensile test was conducted on the JIS No. 5 test piece before processing, so as to
measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile
test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa
as measured by the method described previously, those which showed elongation of 10%
or more were evaluated as high in elongation.
[0178] Evaluation of hydrogen embrittlement resisting property
In order to evaluate the hydrogen embrittlement resisting property, flat test piece
1.2 mm in thickness was subjected to slow stretching rate test (SSRT) with a stretching
speed of 1x10
-4/sec, to determine hydrogen embrittlement risk index (%) defined by the equation shown
below.
Hydrogen embrittlement risk index (%) = 100 × (1-E1/E0)
[0179] E0 represents the elongation before rupture of a steel test piece that does not substantially
contain hydrogen, E1 represents the elongation before rupture of a steel test piece
that has been charged with hydrogen electrochemically in sulfuric acid. Hydrogen charging
was carried out by immersing the steel test piece in a mixed solution of H
2SO
4 (0.5 mol/L) and KSCN (0.01 mol/L) and supplying constant current (100A/m
2) at room temperature.
[0180] A steel sheet having hydrogen embrittlement risk index higher than 50% is likely
to undergo hydrogen embrittlement during use. In the present invention, steel sheets
having hydrogen embrittlement risk index not higher than 50% were evaluated to have
high hydrogen embrittlement resisting property.
[0181] Results of the test are shown in Table 10.
[0182]

[0183]

[0184] The results shown in Tables 9 and 10 can be interpreted as follows (numbers in the
following description are test Nos. in Table 10).
[0185] Test pieces Nos. 301 through 309 (inventive steel sheets 2) and test pieces Nos.
313 through 317 (inventive steel sheets 1) that satisfy the requirements of the present
invention have high strength of 1180 MPa or higher, and show high hydrogen embrittlement
resisting property in harsh environment after the forming process. They also have
high elongation property required of the TRIP steel sheet, thus providing steel sheets
best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere.
[0186] Test pieces Nos. 310 through 312 and 318 that do not satisfy the requirements of
the present invention, in contrast, have the following drawbacks.
[0187] No. 310 made of steel species J-3 that includes excessive C content has carbide precipitated
and residual austenite of longer mean length of minor axis, thus resulting poor performance
in both workability and hydrogen embrittlement resisting property after processing.
[0188] No. 311, martensite steel that is a conventional high strength steel made of steel
species K-3 that includes insufficient Si content, hardly contains residual austenite
and is inferior in hydrogen embrittlement resisting property. It also does not show
the elongation property required of a thin steel sheet.
[0189] No. 312, that was made of a steel that has the composition specified in the present
invention but was not manufactured under the recommended conditions, resulted in the
conventional TRIP steel. As a result, the residual austenite does not have the mean
axis ratio and the mean length of minor axis specified in the present invention, while
the matrix phase is not formed in binary phase structure of bainitic ferrite and martensite,
thus resulting in low strength and poor hydrogen embrittlement resisting property.
[0190] No. 318 includes Al content higher than that specified for the inventive steel sheet
1. As a result, although the predetermined amount of residual austenite is retained,
the residual austenite does not have the mean axis ratio specified in the present
invention, the desired matrix phase is not obtained and inclusions such as AlN are
generated thus resulting in poor hydrogen embrittlement resisting property.
[0191] Then parts were made by using steel species A-3, G-3 shown in Table 9 and comparative
steel sheet (590 MPa class high strength steel sheet of the prior art). Performance
(pressure collapse resistance and impact resistance) of the formed test piece were
studied by conducting pressure collapse test and impact resistance test as follows.
[0192] Pressure collapse test
Maximum tolerable load was determined similarly to Example 1 by using steel species
A-3, G-3 shown in Table 9 and the comparative steel sheet. Absorbed energy was determined
from the area under the load-deformation curve. The results are shown in Table 11.
[0193]
Table 11
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Maximum load |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kN) |
(kJ) |
Symbol A-3 |
1221 |
14 |
8.5 |
11.3 |
0.58 |
Symbol G-3 |
1480 |
14 |
10 |
13.8 |
0.69 |
Comparative steel sheet |
613 |
22 |
0 |
5.7 |
0.33 |
[0194] From Table 11, it can be seen that the part (test piece) made from the steel sheet
of the present invention has higher load bearing capability and absorbs greater energy
than a part made of the conventional steel sheet having lower strength, thus showing
high pressure collapse resistance.
[0195] Impact resistance test
The impact resistance test was conducted similarly to Example 1 on the steel sheets
made of steel species A-3, G-3 shown in Table 9 and the comparative steel sheet. The
results are shown in Table 12.
[0196]
Table 12
Steel sheet used |
Evaluation of test piece |
Steel species |
TS |
EL |
Residual γ |
Energy absorbed |
(MPa) |
(%) |
(Area %) |
(kJ) |
Symbol A-3 |
1221 |
14 |
8.5 |
5.72 |
Symbol G-3 |
1480 |
14 |
10 |
6.88 |
Comparative steel sheet |
613 |
22 |
0 |
3.56 |
[0197] From Table 12, it can be seen that the part (test piece) made from the steel sheet
of the present invention absorbs greater energy than a part made of the conventional
steel sheet that has lower strength, thus showing high impact resistance.
[0198] TEM photographs of the test pieces made in this example are shown as reference. Fig.
12 is a photograph of TEM observation (magnification factor 15000) of No. 301 of the
present invention. Fig. 13 is a photograph of TEM observation (magnification factor
60,000) of a portion shown in the photograph of Fig. 12. From Figs. 12, 13, it can
be seen that the high strength thin steel sheet of the present invention contains
fine residual austenite grains (black portion of bar shape in Figs. 12, 13) specified
in the present invention dispersed therein, and that the residual austenite has the
lath shape specified in the present invention. Fig. 14 is a photograph of TEM observation
of No. 313 of a comparative example. From Fig. 14, it can be seen that the high strength
thin steel sheet of No. 313 contains residual austenite (black portion of somewhat
round shape in Fig. 14), although the residual austenite has a block shape that does
not satisfy the requirements of the present invention.