U.S. patent application number 11/030100 was filed with the patent office on 2005-07-14 for ultra-high strength steel sheet having excellent hydrogen embrittlement resistance, and method for manufacturing the same.
This patent application is currently assigned to Kabushiki Kaisha Kobe Seiko Sho(Kobe Steel, Ltd.). Invention is credited to Akamizu, Hiroshi, Ikeda, Shushi, Makii, Koichi, Mukai, Yoichi, Sugimoto, Koichi.
Application Number | 20050150580 11/030100 |
Document ID | / |
Family ID | 34587737 |
Filed Date | 2005-07-14 |
United States Patent
Application |
20050150580 |
Kind Code |
A1 |
Akamizu, Hiroshi ; et
al. |
July 14, 2005 |
Ultra-high strength steel sheet having excellent hydrogen
embrittlement resistance, and method for manufacturing the same
Abstract
The present invention provides an ultra-high strength steel
sheet having excellent hydrogen embrittlement resistance, which
includes: 0.06 to 0.6% of C; 0.5 to 3% of Si+Al; 0.5 to 3% of Mn;
0.15% or lower of P; and 0.02% or lower of S in terms of mass
percentage, and also includes 3% or higher of residual austenite
structure, 30% or higher of bainitic ferrite structure, and
preferably 50% or lower of polygonal ferrite in terms of an areal
ratio to the entire structure, wherein a mean grain size of bainite
blocks is smaller than 20 .mu.m as determined by comparing
observations of the same region of the bainitic ferrite structure
by EBSP (electron back scatter diffraction pattern) and SEM.
Inventors: |
Akamizu, Hiroshi; (Kobe-shi,
JP) ; Ikeda, Shushi; (Kobe-shi, JP) ; Makii,
Koichi; (Kobe-shi, JP) ; Mukai, Yoichi;
(Kakogawa-shi, JP) ; Sugimoto, Koichi; (Ueda-shi,
JP) |
Correspondence
Address: |
OBLON, SPIVAK, MCCLELLAND, MAIER & NEUSTADT, P.C.
1940 DUKE STREET
ALEXANDRIA
VA
22314
US
|
Assignee: |
Kabushiki Kaisha Kobe Seiko
Sho(Kobe Steel, Ltd.)
Kobe-shi
JP
SHINSHU TLO CO., LTD.
Ueda-shi
JP
|
Family ID: |
34587737 |
Appl. No.: |
11/030100 |
Filed: |
January 7, 2005 |
Current U.S.
Class: |
148/654 ;
148/320 |
Current CPC
Class: |
C21D 8/0278 20130101;
C22C 38/04 20130101; C21D 8/0247 20130101; C22C 38/02 20130101;
C21D 2211/005 20130101; C21D 2211/002 20130101; C22C 38/06
20130101; C21D 1/20 20130101 |
Class at
Publication: |
148/654 ;
148/320 |
International
Class: |
C22C 038/00 |
Foreign Application Data
Date |
Code |
Application Number |
Jan 9, 2004 |
JP |
2004-004727 |
Claims
What is claimed is:
1. An ultra-high strength steel sheet having excellent hydrogen
embrittlement resistance, which includes: 0.06 to 0.6% of C; 0.5 to
3% of Si+Al; 0.5 to 3% of Mn; 0.15% or lower of P; and 0.02% or
lower of S in terms of mass percentage, and also includes 3% or
higher of residual austenite structure and 30% or higher of
bainitic ferrite structure in terms of an areal ratio to the entire
structure, wherein a mean grain size of bainite blocks is smaller
than 20 .mu.m as determined by comparing observations of the same
region of the bainitic ferrite structure by EBSP (electron back
scatter diffraction pattern) and SEM.
2. The ultra-high strength steel sheet according to claim 1,
wherein 5% to 50% of polygonal ferrite structure is included in
terms of an areal ratio to the entire structure.
3. The ultra-high strength steel sheet according to claim 1,
wherein 40% or higher of bainitic ferrite structure is included in
terms of an areal ratio to the entire structure.
4. The ultra-high strength steel sheet according to claim 1,
wherein 50% or higher of bainitic ferrite structure is included in
terms of an areal ratio to the entire structure.
5. The ultra-high strength steel sheet according to claim 1,
further comprising at least one of: 1% or lower (higher than 0%) of
Mo; 0.5% or lower (higher than 0%) of Ni; 0.5% or lower (higher
than 0%) of Cu; and 1% or lower (higher than 0%) of Cr in mass
percentage.
6. The ultra-high strength steel sheet according to claim 1,
further comprising at least one of: 0.1% or lower (higher than 0%)
of Ti; 0.1% or lower (higher than 0%) of Nb; and 0.1% or lower
(higher than 0%) of V in mass percentage.
7. The ultra-high strength steel sheet according to claim 1,
further comprising: 0.003% or lower (higher than 0%) of Ca; and/or
0.003% or lower (higher than 0%) of REM in mass percentage.
8. The ultra-high strength steel sheet according to claim 1, that
has a strength of 1180 MPa class or higher.
9. A method for manufacturing the ultra-high strength steel sheet
according to claim 1, which comprises a heat treatment process of:
heating a steel that contains the components described in claim 1
at a temperature in a range from A3 point to (A3 point+20.degree.
C.) for 10 to 600 seconds after rolling the steel, then cooling the
steel at a mean cooling rate of 3.degree. C./s or more to a
temperature not lower than Ms point and not higher than Bs point,
and keeping the steel in this temperature range for 1 to 1800
seconds.
10. The method according to claim 6, wherein the heat treatment
process is included in a molten zinc plating process.
Description
BACKGROUND OF THE INVENTION
[0001] 1. Technical Field
[0002] The present invention relates to an ultra-high strength
steel sheet having strength of at least 1180 MPa class and
excellent hydrogen embrittlement resistance, and a method for
efficiently manufacturing the ultra-high strength steel sheet.
[0003] 2. Background Art
[0004] There are increasing demands for steel sheets, that are
pressed into forms to be used in such applications as automobiles
and industrial machines, to have both high strength and high
ductility at the same time. Recently needs are increasing for
ultra-high strength steel sheets having strength of at least 1180
MPa class. A type of steel sheet that is regarded as promising one
to satisfy these needs is TRIP (transformation induced plasticity)
steel sheet.
[0005] The TRIP steel sheet includes residual austenite structure
and, when processed to deform at a temperature higher than the
martensitic transformation start point (Ms point), undergoes
considerable elongation due to induced transformation of the
residual austenite (.gamma.R) into martensite by the action of
stress. Known examples 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 ferrite 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 Nisshin Steel Technical Journal No. 43
published in 1980), 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).
[0006] In the realm of ultra-high strength of 1180 MPa upward,
however, the TRIP steel sheet is known to suffer a newly emerging
problem of delayed fracture (crack, etc.) 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 boundary, and
makes the steel brittle. This results in decrease in ductility and
in toughness of the metallic material.
[0007] To counter such problems, researches have been recently
conducted on hydrogen embrittlement of TRIP steel (Non-Patent
Documents Nos. 1 and 2). According to reports of these researches,
all of the types of TRIP steel mentioned above show high hydrogen
embrittlement resistance, while the TBF steel has a particularly
high hydrogen storage capacity. Observation of fracture surface of
the TBF steel shows that quasi cleavage fracture due to hydrogen
storage is suppressed. This fact suggests that the TBF steel has
excellent resistance against delayed fracture. The mechanism behind
this property is supposedly that the TBF steel consists of bainitic
ferrite structure and therefore has a high density of dislocations
in the matrix phase, so that much hydrogen is trapped on the
dislocations, resulting in the storage of higher hydrogen than
other types of TRIP steel.
[0008] However, the TBF steels reported in the literatures
described above show delayed fracture characteristic of about 1000
seconds at the most in terms of crack occurrence time measured by
cathode charging test, indicating the need for further improvement
in the characteristic. Moreover, since the heat treatment
conditions reported in the literatures 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.
[0009] [Non-Patent Document 1] Tomohiko HOJO et. al, "Hydrogen
Embrittlement of Ultra-High Strength Low Alloy TRIP Steel (Part 1:
Hydrogen Storage Characteristic and Ductility", Japan Materials
Science Association, Proceedings of 51.sup.st Academic Lecture
Meeting, 2002, Vol. 8, pp17-18.
[0010] [Non-Patent Document 2] Tomohiko HOJO et. al, "Influence of
Austempering Temperature on Hydrogen Embrittlement of Ultra-high
Strength Low Alloy TRIP steel", CAMP-ISIJ, 2003, Vol. 16, p568
SUMMARY OF THE INVENTION
[0011] The present invention has been made with the background
described above, and has an object of providing a novel TRIP steel
sheet having a bainitic ferrite structure as the matrix phase, that
is an ultra-high strength steel sheet having a tensile strength of
1180 MPa or higher and improved hydrogen embrittlement resistance,
while maintaining the high ductility characteristic of the TRIP
steel, and a method for manufacturing the same.
[0012] In order to achieve the object described above, the
ultra-high strength steel sheet of the present invention
(ultra-high strength steel sheet having tensile strength of 1180
MPa or higher, high capability to elongate and high hydrogen
embrittlement resistance) has such a constitution as 0.06 to 0.6%
of C, 0.5 to 3% of Si.+-.Al, 0.5 to 3% of Mn, 0.15% or lower of P
and 0.02% or lower of S are included in terms of mass percentage,
residual austenite structure occupies 3% or higher and bainitic
ferrite structure occupies 30% or higher in an areal ratio to the
entire structure, wherein a mean grain size of bainite blocks is
smaller than 20 .mu.m as determined by comparing observations of
the same region of the bainitic ferrite structure by EBSP (electron
back scatter diffraction pattern) and SEM, and preferably area
occupied by polygonal ferrite structure is in a range from 5% to
50% in terms of an areal ratio to the entire structure.
[0013] A variation of the steel sheet described above that further
includes at least one of 1% or lower by mass (higher than 0%) of
Mo, 0.5% or lower by mass (higher than 0%) of Ni, 0.5% or lower by
mass (higher than 0%) of Cu and 1% or lower by mass (higher than
0%) of Cr; a variation that further includes at least one of 0.1%
or lower by mass (higher than 0%) of Ti, 0.1% or lower by mass
(higher than 0%) of Nb and 0.1% or lower by mass (higher than 0%)
of V; and a variation that further includes 0.003% or lower by mass
(higher than 0%) of Ca and/or 0.003% or lower by mass (higher than
0%) of REM are all preferred embodiments of the present
invention.
[0014] The method of the present invention that solves the problems
described above is a method of manufacturing the ultra-high
strength steel sheet by applying continuous annealing or plating,
including a heat treatment process wherein a steel that contains
the components described above is kept at a temperature in a range
from A3 point to (A3 point+20.degree. C.) for 10 to 600 seconds,
then cooled at a mean cooling rate of 3.degree. C./s or more to a
temperature not lower than Ms point and not higher than Bs point,
and is kept in this temperature range for I to 1800 seconds.
[0015] According to the present invention, the ultra-high strength
steel sheet having tensile strength of 1180 MPa or higher and
improved hydrogen embrittlement resistance can be manufactured with
high productivity.
BRIEF DESCRIPTION OF THE DRAWINGS
[0016] Other objects and advantages of the invention will become
apparent during the following discussion of the accompanying
drawings, wherein:
[0017] FIG. 1 is a SEM photograph (magnification factor 1500) of
No. 2 (example of the present invention) of Example 1,
[0018] FIG. 2 is a photograph (magnification factor 1500) of the
same region as shown in FIG. 1 using EBSP analysis,
[0019] FIG. 3 is a photograph where residual austenite .gamma.R
(FCC phase) is identified in the EBSP analysis photograph of FIG.
2,
[0020] FIG. 4 is a schematic diagram showing a representative
process according to the method of the present invention, and
[0021] FIG. 5 shows EBSP photographs (magnification factor 5000) of
the example (No. 1) of the present invention in Example 2 and the
comparative example (No. 2).
[0022] FIG. 6 is a schematic perspective view of a member for the
crush resistance test in Example 1.
[0023] FIG. 7 is a schematic side view showing the way in which the
crush resistance test is conducted in Example 1.
[0024] FIG. 8 is a schematic perspective view of a member for the
impact resistance test in Example 1.
[0025] FIG. 9 is a sectional view at A-A in FIG. 8.
[0026] FIG. 10 is a schematic side view showing the way in which
the impact resistance test is conducted in Example 1.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0027] For the purpose of further improving the hydrogen
embrittlement resistance in ultra-high strength region of tensile
strength above 1180 MPa, the present inventors took up TBF steel,
that has bainitic ferrite structure as the matrix phase, from among
various types of TRIP steel, and conducted a research. In
particular, since the method described in Non-Patent Document 2 has
such drawbacks as the high heating temperature results in a low
efficiency of production in practical manufacturing process, high
possibility of damaging the furnace and causing decarburization,
the inventors focused their research on the heating temperature. It
was found that growth of austenite grains can be prevented by
controlling the temperature lower than that employed in the
manufacture of the conventional TBF steel, thus making it possible
to form fine bainite blocks that have been impossible to obtain in
the conventional TBF steel, which lead to improved toughness and
hydrogen embrittlement resistance of the steel, thereby completing
the invention.
[0028] The present invention will now be described in detail.
[0029] [Structure]
[0030] Micro structure that characterizes the present invention
most will be first described.
[0031] The ultra-high strength steel sheet of the present invention
includes a residual austenite structure that occupies 3% or higher
and a bainitic ferrite structure that occupies 30% or higher in an
areal ratio to the entire structure and therefore may be
constituted solely from the residual austenite and the bainitic
ferrite, while 50% or lower (including 0%) of polygonal ferrite
structure may be included where mean grain size of bainite blocks
is smaller than 20 .mu.m as determined by comparing the
observations of the same region of the bainitic ferrite structure
by EBSP and SEM.
[0032] Bainitic Ferrite Structure
[0033] Bainitic ferrite structure is hard in nature, as mentioned
previously, and can readily provide high strength. Also because the
density of dislocations is high in the matrix phase, it has such an
advantage that much hydrogen is trapped on the dislocations,
resulting in the storage of higher hydrogen than other types of
TRIP steel. Further, the bainitic ferrite structure has another
advantage of high capability to elongate since fine residual
austenite grains can be grown in the boundary of lath-shaped
bainitic ferrite. In order to make full use of these advantages, it
is controlled so that the bainitic ferrite structure occupies 30%
or more preferably 40% or higher, and more preferably 50% or higher
in terms of an areal ratio to the entire structure. While the upper
limit of the proportion depends on balancing between the bainitic
ferrite structure and other structures and cannot be definitely
specified, it is recommended to set the upper limit to about 95% or
lower, or preferably 93% or lower in case that the polygonal
ferrite structure is not included. When polygonal ferrite structure
is included, it is recommended to set the upper limit to about 92%
or lower, more preferably 90% or lower.
[0034] The areal ratio of the bainitic ferrite structure is
determined by etching the surface of a steel sheet with Nital
etchant and observing an area (about 50 by 50 .mu.m) with SEM
(scanning electron microscope) photograph (magnification factor
1500) in a surface parallel to the surface on which it was rolled
at a depth of one quarter of the thickness.
[0035] A high resolution FE-SEM equipped with an EBSP detector
(Phillips' XL30S-FEG) was used for the SEM observation in the
present invention. Use of this equipment has an advantage that an
area observed with the SEM can be analyzed by the EBSP detector at
the same time. Use of the FE-SEM also enables it to identify the
bainite blocks in the bainitic ferrite (this subject will be taken
up later).
[0036] Mean Grain Size of Bainite Blocks<20 .mu.m
[0037] The steel sheet or plate of the present invention satisfies
the requirement that mean grain size of bainite blocks is smaller
than 20 .mu.m in the bainitic ferrite that is identified by the
method to be described later in the bainitic ferrite structure.
Thus the present invention is most distinctly characterized in that
hydrogen embrittlement resistance of TBF steel is improved by
forming particularly finer bainite blocks among the bainitic
ferrite structure. The desired property cannot be achieved when
such coarse grains are formed as mean grain size of bainite blocks
is larger than 20 .mu.m. It is preferable that mean grain size of
bainite blocks is as small as possible, preferably 18 .mu.m or
lower, and more preferably 16 .mu.m or lower.
[0038] The bainitic ferrite refers to plate-shaped ferrite in lower
structure having higher density of dislocations (which may or may
not have lath-shaped structure), and is clearly distinguished by
SEM observation from polygonal ferrite that has a lower structure
of extremely low density of dislocations, as described below.
[0039] Polygonal ferrite: black polygonal spots seen in SEM
photograph, that do not include residual austenite or martensite
therein.
[0040] Bainitic ferrite: dark gray spots that often cannot be
distinguished from residual austenite or martensite in SEM
photograph.
[0041] The method of identifying bainite block in bainitic ferrite
will now be described below with reference to FIG. 1 and FIG. 2.
These are the results of observing the same area of sample No. 2 of
Example 1 with the FE-SEM equipped with an EBSP detector, of which
FIG. 1 is an SEM photograph (magnification factor 1500) taken by
the method described above, and FIG. 2 is a photograph of cross
section (magnification factor 1500) in the direction of thickness
subjected to EBSP analysis conducted at the same time in the area
observed with SEM. Hardware and software used in the EBSP
detection, measurement and analysis are those of OIM (Orientation
Imaging Microscopy TM) system manufactured by TexSEM Laboratories
Inc. The measuring interval is 0.1 .mu.m.
[0042] As shown in FIG. 1, since polygonal ferrite and bainitic
ferrite can be distinguished by SEM observation, the region
(bainitic ferrite), from which polygonal ferrite identified by SEM
is excluded, can be easily determined among the structure shown in
FIG. 2 obtained by EBSP analysis upon comparison of the SEM
photograph of FIG. 1 and EBSP photograph of FIG. 2.
[0043] In the bainitic ferrite that is determined as described
above, regions having a difference in orientation not lower than
15.degree. in the inclination angle between adjacent structures (in
the present invention, such regions are regarded as having the same
crystal orientation) are color-identified in red, and grain
boundaries (min. 15.degree., max. 180.degree.) are added to the 001
inverse pole figure. The regions that are color-identified as
described above (regions having a difference in orientation not
lower than 15.degree. in the inclination angle) are defined as
bainite block in the present invention. In other words, bainite
block is defined according to the present invention as a region
that is determined by EBSP analysis to have the same crystal
orientation (region having a difference in orientation hot lower
than 15.degree. in the inclination angle) in the bainitic ferrite
identified by SEM, when the same area is subjected to SEM
observation and EBSP analysis.
[0044] The EBSP method will be briefly described here. EBSP is a
method of determining the crystal orientation at the position where
electron beam is incident, by analyzing Kikuchi pattern obtained
from reflected electrons when the electron beam is directed toward
the surface of specimen. Distribution of orientations over the
specimen surface can be determined by measuring the crystal
orientation at predetermined pitches while scanning the specimen
surface with the electron beam. The EBSP observation has such an
advantage that crystal structures of different orientations in the
direction of thickness, that would be regarded as identical when
observed with a conventional optical microscope, can be
distinguished by the color difference. In case the bainite block
defined on the basis of crystal orientation makes an essential
factor as in the present invention, it is necessary to observe the
structure by the EBSP method.
[0045] For the bainite blocks detected as described above, diameter
of a circle that has the same area as the bainite block is
determined. The diameter of the equivalent circle of bainite block
is determined by using the photograph of EBSP analysis with
magnification factor of 5000. Similarly, diameters of the
equivalent circles of all bainite blocks in the measured area
(about 50 m by 50 .mu.m) are measured and the mean value thereof is
defined as the mean grain size of bainite blocks in the present
invention.
[0046] Residual Austenite Structure (.gamma.R)
[0047] Residual austenite is effective in improving the elongation
property. In order to make full use of this property, the areal
ratio of residual austenite is controlled to be 3% or higher
(preferably 5% or higher, and more preferably 7% or higher) of the
entire structure. Since excessive content of residual austenite
makes it difficult to maintain the ultra-high strength, it is
recommended to keep the content within an upper limit of 30%
(preferably 25%).
[0048] According to the present invention, the residual austenite
structure preferably has the form of lath. The form of lath herein
means a shape with mean axis ratio (major axis/minor axis) of 2 or
more (preferably 4 or more, and more preferably 6 or more). Such a
lath-shaped residual austenite not only has the TRIP effect similar
to that of the conventional residual austenite, but also has an
effect of greatly improving the delayed fracture characteristic,
and is very useful. While there is no particular upper limit to the
mean axis ratio, it is preferably within an upper limit of 30, and
more preferably within 20, since the residual austenite is required
to have a certain thickness in order to achieve the effect of
TRIP.
[0049] In order to make full use of the effect of the lath-shaped
residual austenite, the proportion of the lath-shaped residual
austenite in the residual austenite is preferably as large as
possible. Specific value of the proportion is determined in
accordance to balancing with other structures (bainitic ferrite,
polygonal ferrite, etc.) so as to obtain the desired property. For
the purpose of increasing the strength, it is preferable to set the
proportion of the lath-shaped residual austenite to 50% or higher,
more preferably 60% or higher, still more preferably 70% or higher,
further more preferably 80% or higher, and most preferably 85% or
higher. The entire residual austenite may be lath-shaped residual
austenite, but it is recommended to set an upper limit of about 95%
to the proportion in practice, in consideration of the constraints
imposed by the heating facility and cooling facility and other
factors.
[0050] Concentration of C (C.gamma.R) in the residual austenite is
preferably 0.8% or higher. The value of C.gamma.R has great
influence on the TRIP (strain-induced transformation processing)
characteristics, and is effective in improving the elongation
property when it is controlled to 0.8% or higher. The concentration
is preferably 1% or higher, and more preferably 1.2% or higher.
While the concentration of C.gamma.R is preferably as high as
possible, an upper limit of about 1.6% is supposedly imposed by the
practical processing conditions.
[0051] The residual austenite in the present invention refers to a
region that shows FCC phase (face-centered cubic lattice) when
observed by FE-SEM/EBSP method described previously. Specifically,
measurement is made on an area (about 50 by 50 .mu.m, measurement
pitch of 0.1 .mu.m) in a surface parallel to the surface on which
it was rolled at a depth of one quarter of the thickness. In case
that the surface to be measured is exposed by grinding,
electrolytic grinding is employed in order to prevent the residual
austenite from transforming. Then the specimen placed in a lens
barrel of the SEM is irradiated with electron beam by using the
FE-SEM. EBSP image projected on a screen is captured by a high
sensitivity camera (VE-1000-SIT manufactured by Dage-MIT Inc.) and
is imported into a computer. The image is analyzed by the computer,
and compared with a pattern generated by simulation using a known
crystal system (FCC phase (face-centered cubic lattice) in the case
of residual austenite) so as to color-identify the FCC phase. The
areal ratio of the regions thus identified is defined as the areal
ratio of residual austenite. Hardware and software used in the
analysis described above are those of OIM (Orientation Imaging
Microscopy TM) system manufactured by TexSEM Laboratories Inc.
[0052] FIG. 3 shows the EBSP photograph of FIG. 2 wherein FCC phase
is color-identified (magnification factor 1500). In FIG. 3, the
region indicated by an arrow (.rarw.) is the residual austenite
(.gamma.R).
[0053] Polygonal Ferrite
[0054] In the present invention, polygonal ferrite means a ferrite
structure that includes no or very few dislocations.
[0055] In the present invention, polygonal ferrite may be optional
structure and 0% of polygonal structure is included. In order to
make full use of the effect of polygonal ferrite to improve the
elongation property, the areal ratio of polygonal ferrite is
preferably controlled to 5% or higher, and more preferably 10% or
higher. In order to improve the elongation property, the content of
polygonal ferrite is preferably as high as possible. However, since
when the content of polygonal ferrite exceeds 50%, it becomes
difficult to ensure the required level of strength, it is
controlled within an upper limit of 50%, preferably 40% and more
preferably 30%.
[0056] The areal ratio of polygonal ferrite is calculated by the
following equation:
Areal ratio of polygonal ferrite=100-[Areal ratio of bainitic
ferrite (%)]-[Areal ratio of residual austenite
[0057] Note: The areal ratio of bainitic ferrite and the areal
ratio of residual austenite are measured by the method described
previously.
[0058] Others
[0059] The steel sheet of the present invention may be constituted
either from only the structures described above (namely, a
composite structure of bainitic ferrite and residual austenite or a
composite structure of bainitic ferrite, polygonal ferrite and
residual austenite), or may include other structure (for example,
martensite, etc.) to such an extent that the effect of the present
invention is not disturbed. 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 (for example, total areal ratio is within an upper
limit of 10% at the most).
[0060] Now the essential components of the steel sheet of the
present invention will be described. Hereinafter concentrations of
components are all given in terms of mass percentage.
[0061] C: 0.06 to 0.6%
[0062] C is an essential element for ensuring high strength and
maintaining residual austenite. Particularly it is important to
contain a sufficient content of C in the austenite phase, so as to
maintain the desired austenite phase to remain even at the room
temperature, and C content is useful in achieving better compromise
between the strength and the elongation property. Adding 0.25% or
higher of C increases the amount of the residual austenite and also
increases the concentration of C in the residual austenite, thus
enabling it to obtain very high strength and elongation.
[0063] However, the concentration of C higher than 0.6% does not
increase the effect beyond saturation and may also result in
defects due to segregation. The concentration of C higher than
0.25% leads to poor weldability.
[0064] Therefore, when emphasis is placed on weldability, it is
preferable to control the concentration of C in a range from 0.06%
to 0.25% (more preferably within 0.2%, and further more preferably
within 0.15%). For such applications that require high elongation
but does not involve point welding, it is recommended to control
the concentration of C in a range from 0.25% to 0.6% (more
preferably 0.3% or higher).
[0065] Si+Al: 0.5 to 3%
[0066] Si and Al have an effect of suppressing residual austenite
from decomposing and producing carbide. Si, in particular, is also
effective in enhancing solid solution. In order to make full use of
this effect, it is necessary to add Si and Al in total
concentration of 0.5% or more preferably 0.7% or higher, and more
preferably 1% or higher. However, adding Si and Al in total
concentration of higher than 3% does not increase the effect beyond
saturation and cannot be justified economically. Further, too much
content leads to hot rolling brittleness. Therefore, the
concentration is controlled within an upper limit of 3%, preferably
within 2.5% and more preferably within 2%.
[0067] Mn: 0.5 to 3%
[0068] 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 0.5% or more
preferably 0.7% or higher, and more preferably 1% or higher.
However, adding Mn in concentration higher than 3% causes adverse
effects. The concentration is preferably within 2.5% and more
preferably within 2%.
[0069] P: 0.15% or Lower (Gigher Than 0%)
[0070] P is an element that is effective in obtaining desired
residual austenite. In order to make full use of this effect, it is
recommended to add P in concentration of 0.03% or higher
(preferably 0.05% or higher). However, adding P in concentration
higher than 0.15% adversely affects the ease of secondary
processing. Thus the concentration is more preferably within
0.1%.
[0071] S: 0.02% or Lower (Higher Than 0%)
[0072] S forms sulfide inclusion such as MnS that initiates crack
and adversely affects the workability of the steel. Concentration
of S is preferably within 0.02% and more preferably within 0.015%.
Effect of decreasing the S content in suppressing the deterioration
of the workability does not hold when the concentration of S
decreases below 0.003%. Decreasing the S content below this level
is expensive without providing benefit. Therefore, it is
recommended to set the lower limit of S content to over 0.003%, or
preferably to 0.005% or higher.
[0073] While the steel of the present invention includes the
elements described above as the fundamental components with the
rest substantially consisting of iron and impurities, the following
elements may be added to such an extent that does not disturb or
compromise the effect of the present invention.
[0074] At Least One of Mo: 1% or Lower (Higher Than 0%), Ni: 0.5%
or Lower (Higher Than 0%), Cu: 0.5% or Lower (Higher Than 0%) and
Cr: 1% or Lower (Higher Than 0%)
[0075] These elements are effective in strengthening the steel and
stabilizing and ensuring the predetermined amount of residual
austenite. In order to make full use of this effect, it is
recommended to add Mo in concentration of 0.05% or higher
(preferably 0.1% or higher), Ni in concentration of 0.05% or higher
(preferably 0.1% or higher), Cu in concentration of 0.05% or higher
(preferably 0.1% or higher) and Cr in concentration of 0.05% or
higher (preferably 0.1% or higher). However, the effects described
above reach saturation when higher than 1% of Mo and Cr, and higher
than 0.5% of Ni and Cu are added, resulting in economical
disadvantage. It is more preferable to add 0.8% or lower Mo, 0.4%
or lower Ni, 0.4% or lower Cu and 0.8% or lower Cr.
[0076] At Least One of Ti: 0.1% or Lower (Higher Than 0%), Nb: 0.1%
or Lower (Higher Than 0%) and V: 0.1% or Lower (Higher Than 0%)
[0077] These elements have the effects of enhancing precipitation
and making the structure finer, and are effective in strengthening
the steel. In order to make full use of these effects, it is
recommended to add Ti in concentration of 0.01% or higher
(preferably 0.02% or higher), Nb in concentration of 0.01% or
higher (preferably 0.02% or higher) and V in concentration of 0.01%
or higher (preferably 0.02% or higher). However, the effects
described above reach saturation when the concentration of any of
these elements exceeds 0.1%, resulting in economical disadvantage.
It is more preferable to add 0.08% or lower Ti, 0.08% or lower Nb
and 0.08% or lower V.
[0078] Ca: 0.003% or Lower and/or REM: 0.003% or Lower (Higher Than
0%)
[0079] Ca and REM (rare earth element) are effective in controlling
the form of sulfide in the steel and improve the workability of the
steel. Sc, Y, La and the like may be used as the rare earth element
in the present invention. In order to achieve the effect described
above, it is recommended to add each of these elements in
concentration of 0.0003% or higher (preferably 0.0005% or higher),
However, the effects described above reach saturation when the
concentration exceeds 0.003%, resulting in economical disadvantage.
It is more preferable to keep the concentration within 0.0025%.
[0080] The method for manufacturing the steel sheet of the present
invention will now be described.
[0081] The method of the present invention is characterized in that
the steel that has the composition described above is kept at a
temperature in a range from A3 point to (A3 point+20.degree. C.)
for 10 to 600 seconds, then cooled at a mean cooling rate of
3.degree. C./s or more to a temperature not lower than Ms point and
not higher than Bs point, and is kept in this temperature range for
1 to 1800 seconds. Each stage of this process will be described
below with reference to FIG. 4 that schematically illustrates the
method of the invention.
[0082] First, the steel that has the composition described above is
heated to a temperature in a range from A3 point to (A3
point+20.degree.C.) (T1 in FIG. 4) and is kept at this temperature
for 10 to 600 seconds (t1 in FIG. 4). T1 (soaking temperature) and
t1 (soaking time) are very important for obtaining the desired
bainite blocks. Austenite grains will grow to produce coarse
bainite blocks, when T1 is higher than (A3 point+20.degree. C.) or
t1 is longer than 600 seconds.
[0083] When T1 is lower than the temperature of A3 point, on the
other hand, predetermined bainitic ferrite structure cannot be
obtained. When t1 is lower than 10 seconds, austenizing process
does not proceed sufficiently, leaving cementite and other carbides
to remain.
[0084] Based on these considerations, it is preferable to set T1
(soaking temperature) in a range from 650 to 900.degree. C. and t1
(soaking time) in a range from 30 to 300 seconds, and more
preferably in a range from 60 to 240 seconds.
[0085] Then the steel is cooled at a mean cooling rate of 3.degree.
C./s (CR1 in FIG. 4) or higher to a temperature (T2 in FIG. 1) not
lower than Ms point and not higher than Bs point, and is kept in
this temperature range for 1 to 1800 seconds (t2 in FIG. 4).
[0086] This process is designed for the purpose of forming the
desired bainitic ferrite structure (which may include polygonal
ferrite structure) and avoiding the formation of pearlite structure
that is not desirable for the present invention (to suppress the
areal ratio of pearlite within 10% at the most).
[0087] The purpose of cooling the heated steel at the mean cooling
rate of 3.degree. C./s (CR1) or more is to avoid the pearlite
transformation zone so as to prevent the generation of pearlite
structure. The mean cooling rate is desired to be as fast as
possible, preferably 10.degree. C. per second or more (more
preferably 20.degree. C. per second or more). While the steel may
be cooled quickly to the predetermined temperature T2 (1-stage
cooling) as shown in FIG. 1, it is difficult to form the polygonal
ferrite structure in 1-stage cooling process. Therefore, when it is
desired to form the polygonal ferrite structure as well, it is
recommended to employ multi-stage cooling process by dividing the
cooling process into several stages. In this case, too, it is
recommended to set the mean cooling rate in each cooling stage to
3.degree. C. per second or higher (preferably 10.degree. C. per
second or higher and more preferably 20.degree. C. per second or
higher).
[0088] After quenching to the temperature (T2) that is not lower
than Ms point and not higher than Bs point, the steel is kept at
this temperature so as to undergo isothermal transformation in
which the desired bainitic ferrite structure is formed. When the
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 areal ratio of austenite
decreases.
[0089] When the temperature holding period t2 is longer than 1800
seconds, density of dislocations in bainitic ferrite becomes low
and the desired residual austenite cannot be obtained. When t2 is
lower than 1 second, on the other hand, desired bainitc ferrite
cannot be obtained. Length of t2 is preferably from 30 to 1200
seconds, and more preferably from 60 to 600 seconds.
[0090] 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 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. Further, the plating may also be alloyed.
[0091] 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 A3 point, the steel sheet is cooled at a mean cooling rate of
about 30.degree. C./s and is wound up at a temperature from about
500 to 600.degree. C. In case that the hot rolled steel sheet has
poor 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 30%. Cold rolling beyond 30% leads to excessive
rolling load that makes it difficult to carry out the cold
rolling.
[0092] A steel sheet according to the present invention has not
only a high capability to elongate and a high hydrogen
embrittlement resistance, but also excellent properties for safety
in collision. Therefore, the steel sheet is used, for example, for
construction members of a vehicle and industrial machinery.
Especially, the steel sheet is suitable for automobile members such
as crush members, for example, side members of the front portion
and the back portion and crush boxes, construction members, for
example, pillar members (such as center pillar reinforce), roof
rail reinforces, side sills, floor members and kick portions, and
impact absorbing members, for example, bumper reinforces and door
impact beams.
[0093] Now the present invention will be described in detail below
by way of examples. It is understood, however, that the present
invention is not limited by these examples, and various
modifications that do not deviate from the spirit of the present
invention described herein are all within the scope of the present
invention.
EXAMPLES
Example 1
Investigation On Composition
[0094] In this example, steel specimens A through P having the
compositions shown in Table 1 (rest of the composition consists of
iron and impurities, and the concentrations in the table being
given in mass percentage) was made by vacuum melting to obtain an
experimental slab that was subjected to the process described below
(hot rolling.fwdarw.cold rolling.fwdarw.continuous annealing) to
turn into a hot rolled steel sheet having thickness of 3.2 mm
(thickness:2.5mm in No. Q to T), that was then pickled to remove
scales from the surface and was cold rolled to a thickness of 1.2
mm.
[0095] Hot rolling process: Starting temperature (SRT) 1150.degree.
C., finishing temperature (FDT) 850.degree. C., cooling rate of
40.degree. C./s and take-up temperature 550.degree. C.
[0096] Cold rolling process: Rolling ratio 50%.
[0097] Continuous annealing process: Each steel specimen was kept
at a temperature (A3 point+15.degree. C.) (T1 in Table 1) for 120
seconds (t1 in Table 1), then cooled (water cooling) at a mean
cooling rate of 25.degree. C./s (CR1 in Table 1) to 420.degree. C.
(T2 in Table 1), and was then kept at 420.degree. C. for 120
seconds (t2 in Table 1).
[0098] Tensile strength (TS), elongation (total elongation E1) and
hydrogen embrittlement resistance (cathode CH life) were measured
on each of the steel sheets obtained as described above. The areal
ratio of structure and mean grain size of bainite blocks in each
steel sheet were measured by the methods described previously.
[0099] [Measurement of Tensile Strength (TS) and Elongation]
[0100] Tensile test was conducted by using JIS No. 5 test piece to
measure tensile strength (TS) and elongation (E1). Speed of
elongation in the tensile test was set to 1 mm/sec. in the present
invention. The steel sheet of the present invention is one that
shows tensile strength of 1180 MPa or higher in the test described
above, and those that show 13% or higher elongation are evaluated
as excellent in elongation property in the present invention.
[0101] [Measurement of Hydrogen Embrittlement Resistance]
[0102] Hydrogen embrittlement resistance was measured by using
rectangular test pieces of the steel sheets described above having
a size of 15 mm by 65 mm. The rectangular test piece was loaded
with a pressure of 980 MPa by 4-point bending and was subjected to
a potential of -80 mV, a potential baser than the natural
potential, using a potentiostat in a solution of 0.5 mol of
sulfuric acid and 0.01 mol of KSCN, and the elapsed time before
crack occurred was measured thereby to evaluate the hydrogen
embrittlement resistance (cathode CH life). In the present
invention, those that survived for 1000 seconds or more without
crack are evaluated as excellent in hydrogen embrittlement
resistance.
[0103] The test results are shown in Table 2.
1 TABLE 1 A3 trans- for- Si + mation C Si AI Al Mn P S Others point
A 0.03 1.5 0.03 1.53 1.5 0.02 0.005 -- 871 B 0.3 1.5 0.03 1.53 1.5
0.02 0.005 -- 795 C 0.5 1.5 0.03 1.53 1.5 0.02 0.005 -- 763 D 0.3
0.5 0.5 1.00 1.5 0.02 0.005 -- 652 E 0.3 0.1 0.03 0.13 1.5 0.02
0.005 -- 732 F 0.3 1.5 0.03 1.53 0.3 0.02 0.005 -- 831 G 0.3 1.5
0.03 1.53 5.0 0.02 0.005 -- 690 H 0.3 1.5 0.03 1.53 1.5 0.02 0.005
Mo: 0.2 801 I 0.3 1.5 0.03 1.53 1.5 0.02 0.005 Ni: 0.2 792 J 0.3
1.5 0.03 1.53 1.5 0.02 0.005 Cu: 0.2 799 K 0.3 1.5 0.03 1.53 1.5
0.02 0.005 Cr: 0.2 780 L 0.3 1.5 0.03 1.53 1.5 0.02 0.005 Ti: 0.03
783 M 0.3 1.5 0.03 1.53 1.5 0.02 0.005 Nb: 0.03 795 N 0.3 1.5 0.03
1.53 1.5 0.02 0.005 V: 0.03 798 O 0.3 1.5 0.03 1.53 1.5 0.02 0.005
Ca: 0.001 794 P 0.08 1.5 0.03 1.53 1.5 0.02 0.005 Mo: 0.2 843 Ti:
0.03 Q 0.2 1.5 0.03 1.53 2.5 0.01 0.005 Nb: 0.05 798 Mo: 0.2 R 0.2
1.5 0.03 1.53 3.0 0.01 0.004 Nb: 0.05 783 Mo: 0.2 S 0.17 1.5 0.03
1.53 2.5 0.01 0.005 Nb: 0.05 788 Mo: 0.2 Ti: 0.05 Cu: 0.3 Ni: 0.2 T
0.17 1.35 0.04 1.37 2.0 0.01 0.005 -- 810
[0104]
2 TABLE 2 Continuous annealing process Structure Properties T1 t1
CR1 T2 t2 Areal ratio (%) Bainite TS EL Cathode CH No. Type
(.degree. C.) (S) (.degree. C./s) (.degree. C.) (s) Bainite Ferrite
Residual .gamma. block (.mu.m) (MPa) (%) life (Seconds) 1 A 886 120
25 420 120 16 79 5 10 690 30 -- 2 B 810 120 25 420 120 82 10 8 11
1198 15 1519 3 C 778 120 25 420 120 80 10 10 12 1230 15 1424 4 D
667 120 25 420 120 77 15 8 13 1220 15 1529 5 E 747 120 25 420 120
90 10 0 12 1200 10 1374 6 F 846 120 25 420 120 88 11 1 11 1180 12
1477 7 G 705 Process interrupted -- -- -- -- -- -- -- by hot
rolling crack. 8 H 816 120 25 420 120 82 10 8 9 1230 15 1552 9 I
807 120 25 420 120 82 10 8 9 1232 15 1647 10 J 814 120 25 420 120
82 10 8 8 1234 15 1641 11 K 795 120 25 420 120 82 10 8 8 1236 15
1599 12 L 798 120 25 420 120 82 10 8 10 1238 15 1629 13 M 810 120
25 420 120 82 10 8 10 1240 14 1619 14 N 813 120 25 420 120 82 10 8
10 1242 14 1613 15 O 809 120 25 420 120 82 10 8 11 1241 14 1587 16
P 858 120 25 420 120 82 10 8 9 1190 15 1525 17 Q 815 120 20 400 120
92 2 6 6 1510 10 1215 18 R 800 120 20 400 120 90 2 8 5 1530 10 1220
19 S 805 120 20 300 120 89 4 7 8 1410 12 1310 20 T 880 120 WQ 200
120 0 35 1 -- 1045 16 -- Residual .gamma.: Residual austenite
[0105] These results can be interpreted as follows (all the numbers
in the following discussion are the experiment numbers shown in
Table 2).
[0106] Nos. 2 through 4 and 8 through 16 are all examples of the
present invention manufactured in accordance to the method of the
present invention using the steel of types that satisfy the
conditions of the present invention (B through D and H through P in
Table 1), and show excellent performance in both elongation
property and hydrogen embrittlement resistance in the realm of
ultra-high strength of 1180 MPa upward.
[0107] The steel of types that do not satisfy the conditions of the
present invention (A and E through G in Table 1), on the other
hand, have such drawbacks as described below.
[0108] No. 1 is an example made of a steel of type A having lower C
content, where the predetermined amount of the bainitic ferrite
structure (hard structure) could not be formed while excessive
ferrite structure was observed, showing lower strength. Hydrogen
embrittlement resistance was not measured because of such a low
strength that load could not be applied in 4-point bending
test.
[0109] No. 5 is an example made of a steel of type E that has a low
total concentration of Si and Al, where the desired residual
austenite is not obtained and elongation is low.
[0110] No. 6 is an example made of a steel of type F that has a low
concentration of Mn, where the desired residual austenite is not
obtained and, as a result, elongation property is low.
[0111] No. 7 is an example made of a steel of type G that has a
high concentration of Mn, where excessively high strength caused
rolling crack during hot rolling, thus making it impossible to
carry out the subsequent annealing process.
[0112] No. 20 is a conventional dual-phase steel sheet, and has a
lower strength because it does not contain bainite. Hydrogen
embrittlement resistance was not measured because it did not have a
sufficient strength.
[0113] Next, formed products made of steel sheet No. 17 or No. 20
were evaluated in crush resistance, impact resistance and hydrogen
embrittlement resistance in order to examine the properties as
formed product.
[0114] <Crush Resistance Test>
[0115] A member 1 as shown in FIG. 6 (hat channel member) was made
of No. 17 or No. 19. A crush resistance test for the member was
conducted in the following way. Spot welding was performed in 35 mm
pitch for the spot welding positions 2 in the member 1, as shown in
FIG. 6, wherein an electrode of 6 mm diameter was used and a
current 0.5 kA lower than the splash current was applied. Then, as
shown in FIG. 7, a metal mold was pushed from above onto the center
of the member 1 in the longitudinal direction and the maximum load
was obtained. At the same time, absorbed energy was obtained
according to the area in load displacement diagram. The results are
shown in Table 4.
3 TABLE 3 Used Steel Sheet Test Results Retained .gamma. Maximum
Absorbed Steel TS EL (areal Load Energy No. Composition (MPa) (%)
ratio %) (kN) (kJ) 17 Q 1510 10 6 13.6 0.61 20 T 1045 16 0.2 10.2
0.48
[0116] Table 3 shows that the member made of No. 17 steel sheet has
a higher load and higher energy absorption property than one made
of a conventional low strength steel sheet and that it has an
excellent crush resistance.
[0117] <Impact Resistance Test>
[0118] A member 4 as shown in FIG. 8 (hat channel member) was made
of No. 17 or No. 19. An impact resistance test for the member was
conducted in the following way. FIG. 9 is a sectional view of the
member 4 at A-A in FIG. 8. Spot welding was performed for the spot
welding positions 5 in the member 4. Then, as schematically shown
in FIG. 10, the member 4 was installed on a base 7 and a hammer 6
(110 kg mass) was dropped from the position 11 m high above the
member 4. Absorbed energy until the member was deformed (in the
height direction) by 40 mm was obtained. The results are shown in
Table 4.
4 TABLE 4 Used Steel Sheet Retained Test Results Steel TS EL
.gamma. (areal Absorbed No. Composition (MPa) (%) ratio %) Energy
(kJ) 17 Q 1510 10 6 6.23 20 T 1045 16 0.2 4.38
[0119] Table 4 shows that the member made of No. 17 steel sheet has
an energy absorption property higher than one made of a
conventional low strength steel sheet and that it has excellent
impact resistance.
[0120] <Evaluation of Hydrogen Embrittlement Resistance>
[0121] Members for real use were formed using Nos. 17 and 19 steel
sheets. Hydrogen embrittlement resistance was evaluated in the
state of formed products. Specifically, Nos. 17 and 19 steel sheets
were press-worked into center pillar reinforces, door impact beams
and roof rail reinforces and then dipped into 5% hydrochloric acid.
It was examined whether cracks generated in the test members or
not. The results are shown 15 in Table 5.
5TABLE 5 Crack generation in 24 Steel hours after dipping into No.
Composition Formed Member acid solution 17 Q center pillar No door
impact beam No roof rail No 20 T center pillar No door impact beam
No roof rail No
[0122] Table 5 shows that the member made of No. 17 steel sheet
suffers no crack generation and an excellent hydrogen embrittlement
resistance in spite of its high strength.
Example 2
Investigation of Manufacturing Conditions
[0123] In this example, an experimental slab made by using the
steel of type B shown in Table 1 (steel that satisfies the
conditions of the present invention) was subjected to hot rolling
and cold rolling under the same conditions as those of the Example
1, followed by continuous annealing under various conditions shown
in Table 6, thereby to obtain the cold rolled steel sheets Nos.1
through 11, all 1.2 mm in thickness.
[0124] Then structures and various properties of these steel sheets
were investigated similarly to the Example 1. The results are shown
in Table 6.
6 TABLE 6 Continuous annealing or plating process Structure
Properties t1 CR1 T2 t2 Zn-GA Areal ratio (%) Bainite TS EL Cathode
CH No. T1 (.degree. C.) (S) (.degree. C./s) (.degree. C.) (s)
(.degree. C.) Bainite Ferrite Residual .gamma. Other block (.mu.m)
(MPa) (%) life (Seconds) 1 810 120 25 420 120 -- 82 10 8 0 10 1198
15 1519 2 930 120 25 420 120 -- 82 10 8 0 30 1198 15 1119 3 760 120
25 420 120 -- 10 75 15 0 5 980 18 -- 4 810 300 25 420 120 -- 82 10
8 0 25 1198 22 1185 5 810 1 25 420 120 -- 89 10 1 0 11 1198 11 1165
6 810 120 1 420 120 -- 0 40 0 P(60) -- 1010 11 819 7 810 120 25 600
120 -- 20 10 1 P(69) 11 1200 11 898 8 810 120 25 200 120 -- 89 10 1
0 12 1220 10 1165 9 810 120 25 420 0.5 -- 0 10 8 M(82) 11 1250 8
780 10 810 120 25 420 3600 -- 89 10 1 0 11 1182 11 919 11 810 120
25 420 120 550 82 10 8 0 10 1199 15 1519 P: Pearlite M: Martensite
Residual .gamma.: Residual austenite
[0125] No. 1 is an example of the present invention manufactured in
accordance to the method of the present invention. No. 11 is an
example of the present invention that was subjected, after the
process described above, to alloying treatment (dipped in molten
zinc and then heat-treated at 500.degree. C. for alloying). Both
specimens have ultra-high strength above 1180 MPa, and are
excellent in both elongation and hydrogen embrittlement
resistance.
[0126] Specimens Nos.2 through 10 that do not satisfy some of the
conditions of the present invention, on the other hand, have such
drawbacks as described below.
[0127] No. 2, that was heated at temperature T1 of 930.degree. C.,
higher than the upper limit of the present invention (815.degree.
C. since A3 point of steel type B is 795.degree. C.), included
coarse bainite blocks and showed lower hydrogen embrittlement
resistance.
[0128] No. 3, that was heated at temperature T1 of 760.degree. C.,
lower than the lower limit of the present invention (A3 point of
steel type B=795.degree. C.), did not include predetermined amount
of bainitic ferrite structure, showing lower strength. Hydrogen
embrittlement resistance was not measured because of such a low
strength that load could not be applied in 4-point bending
test.
[0129] No. 4, that was heated for a long heating time t1, resulted
in the growth of austenite grains and coarse bainite blocks, with
lower hydrogen embrittlement resistance.
[0130] In No. 5, that was heated for a short heating time t1,
austenizing did not proceed sufficiently thus leaving cementite to
remain, with such a result that the desired residual austenite was
not obtained and elongation was low, while hydrogen embrittlement
resistance decreased.
[0131] In No. 6, that was cooled at a slow cooling rate CR1, much
pearlite structure was formed and the predetermined amount of
bainitic ferrite was not obtained, thus desired elongation and
hydrogen embrittlement resistance could not be achieved. In
addition, predetermined amount of residual austenite was not
obtained and elongation was low.
[0132] In No. 7, that was heated to beyond Ms point at heating
temperature T2 of 600.degree. C. after cooling, much pearlite
structure was formed and the predetermined amount of bainitic
ferrite was not obtained, thus desired strength could not be
achieved and hydrogen embrittlement resistance deteriorated. In
addition, predetermined amount of residual austenite was not
obtained and elongation was low.
[0133] In No. 8, that was heated to below Bs point at heating
temperature T2 of 200.degree. C. after cooling, the predetermined
amount of residual austenite was not obtained, with elongation and
hydrogen embrittlement resistance becoming lower.
[0134] In No. 9, that was heated for a short heating time t2 after
cooling, predetermined bainitic ferrite was not obtained, resulting
in low hydrogen embrittlement resistance. In addition, martensite
was formed and elongation became lower.
[0135] In No. 10, that was heated for a long heating time t2 after
cooling, decomposition of residual austenite proceeded, thus
predetermined amount of residual austenite was not obtained and
elongation was low.
[0136] FIG. 5 shows EBSP photographs (color-identified with
magnification factor 5000) of the example (No. 1) of the present
invention and the comparative example (No. 2) for reference. From
FIG. 5, it can be seen that No. 1 made by the method of the present
invention shows the desired fine bainite blocks formed therein,
while the comparative example No. 2 that does not fall in the scope
of the present invention shows coarse bainite blocks formed
therein.
* * * * *