DESCRIPTION
Title of Invention: High Strength Steel Sheet and High
Strength Galvanized Steel Sheet Excellent in Shapeability
and Methods of Production of Same
Technical Field
[OOOl] The present invention relates to high strength
steel sheet and high strength galvanized steel sheet
which are excellent in shapeability and to methods of
production of the same.
Background Art
[0002] In recent years, there have been increasing
demands for higher strength in the steel sheet which is
used for automobiles etc. In particular, for the purpose
of improving collision safety etc., high strength steel
sheet with a tensile maximum stress of 900 MPa or more is
also being used. Such high strength steel sheet is
inexpensively formed in large volumes by press working it
in the same way as soft steel sheet and is being used as
structural members.
[0003] However, in recent years, along with the rapid
increase in strength of high strength steel sheet, in
particular in high strength steel sheet with a tensile
maximum stress of 900 MPa or more, the problem has arisen
of the shapeability becoming insufficient and of working
accompanied with local deformation such as stretchformability
becoming difficult. For this reason, in high
strength steel sheet with a high tensile maximum stress
as well, realization of sufficient workability has become
demanded.
[0004] PLT 1 discloses, as art for improving the
bendability of high strength steel sheet, steel sheet
with a tensile strength of 780 to 1470 MPa, a good shape,
and excellent bendability which is obtained by taking
steel sheet which has a microstructure mainly comprised
of bainite or tempered martensite, making the amount of
Si which is contained in the steel, by mass%, 0.6% or
less, cooling down to a temperature at least 50°C lower
than a predetermined bainite transformation temperature
to promote transformation from austenite to bainite or
martensite and thereby rendering the volume rate of
residual austenite which is contained in the structure
and has a martensite transformation point of -196°C or
more 2% or less.
[OOOS] PLT 2 discloses, as art for improving the
shapeability of high strength steel sheet, the method of
improving the ductility and stretch flangeability by
cooling steel sheet which has been hot rolled down to
500°C or less, coiling it, then reheating it to 550 to
700°C, then successively performing a cold rolling process
and continuous annealing process so that a second phase
which contains residual austenite and further contains a
low temperature transformation phase becomes fine in
average particle size and so that the amount of residual
austenite, amount of solid solution C in the residual
austenite, and average particle size satisfy
predetermined relationship formulas.
[0006] PLT 3 discloses, as art for improving the
stretch flangeability of high strength steel sheet, steel
sheet which is reduced in standard difference in hardness
inside of the steel sheet and which is given equivalent
hardness in the entire steel sheet region.
[0007] PLT 4 discloses, as art for improving the
stretch flangeability of high strength steel sheet, steel
sheet which is reduced in hardness of hard portions by
heat treatment and which is reduced in hardness
difference with the soft parts.
[OOOS] PLT 5 discloses, as art for improving the
stretch flangeability of high strength steel sheet,
rendering the hard portions the relatively soft bainite
so as to reduce the difference in hardness from soft
parts.
[0009] PLT 6 discloses, as art for improving the
stretch flangeability of high strength steel sheet, steel
sheet which has a structure comprised of, by area rate,
40 to 70% of tempered martensite and a balance of ferrite
where a ratio between an upper limit value and a lower
limit value of a concentration of Mn in the cross-section
of the thickness direction of the steel sheet is reduced.
Citations List
Patent Literature
[OOlO] PLT 1: Japanese Patent Publication No. 10-
280090A
PLT 2: Japanese Patent Publication No. 2003-183775A
PLT 3: Japanese Patent Publication No. 2008-266779A
PLT 4: Japanese Patent Publication No. 2007-302918A
PLT 5: Japanese Patent Publication No. 2004-263270A
PLT 6: Japanese Patent Publication No. 2010-65307A
Summary of Invention
Technical Problem
[OOll] In the high strength steel sheet which is
described in PLT 1, there is the problem that in the
steel sheet structure, there is little ferrite and
residual austenite for improving the ductility and
therefore sufficient ductility cannot be obtained.
[0012] The method of production of high strength steel
sheet according to PLT 2 requires a large scale reheating
apparatus, so there is the problem that the manufacturing
cost increases.
[0013] In the arts which are described in PLTs 3 to 6
as well, the workability in high strength steel sheet
with a tensile maximum strength of 900 MPa or more is
insufficient.
[OOlS] The present invention was made in consideration
of the above problems and has as its object the provision
of high strength steel sheet and high strength galvanized
steel sheet which are excellent in shapeability and
methods of production of the same by which a tensile
maximum strength 900 MPa or more high strength is secured
while excellent ductility and stretch flangeability are
obtained.
Solution to Problem
[0015] The inventors etc. engaged in intensive studies
on the steel sheet structure and method of production for
obtaining excellent ductility and stretch flangeability
in high strength steel sheet. As a result, they
discovered that by making the steel ingredients suitable
ranges and further by establishing suitable annealing
conditions after cold rolling, it is possible to make the
ratio of the residual austenite phase in the steel sheet
structure a predetermined range while lowering the
martensite transformation start temperature of the
residual austenite phase and that by producing high
strength steel sheet under such conditions and
controlling the ratio of the residual austenite phase in
the steel sheet structure and the martensite
transformation point to suitable ranges, a 900 MPa or
higher tensile maximum strength is secured while the
ductility and stretch flangeability (hole expandability)
are improved and excellent shapeability is obtained.
[0016] The present invention was made as a result of
further studies based on the above findings and has as
its gist the following:
[0017] (1) High strength steel sheet which is
excellent in shapeability which contains, by mass%, C:
0.075 to 0.300%, Si: 0.70 to 2.50%, Mn: 1.30 to 3.50%, P:
0.001 to 0.030%, S: 0.0001 to 0.0100%, Al: 0.005 to
1.500%, N: 0.0001 to 0.0100%, and 0: 0.0001 to 0.0100%,
which contains, as optional elements, one or more of Ti:
0.005 to 0.150%, Nb: 0.005 to 0.150%, B: 0.0001 to
0.0100%, Cr: 0.01 to 2.00%, Ni: 0.01 to 2.00%, Cu: 0.01
to 2.00%, Mo: 0.01 to 1.00%, V: 0.005 to 0.150%, and one
or more of Ca, Ce, Mg, Zr, Hf, and REM: total 0.0001 to
0.5000%, and has a balance of iron and unavoidable
impurities, wherein the structure of the steel sheet
contains, by volume fraction, 2 to 20% of residual
austenite phase, and the residual austenite phase has a
martensite transformation point of -60°C or less.
[0018] (2) The high strength steel sheet which is
excellent in shapeability according to (I), characterized
in that a ratio of the residual austenite phase which
transforms to martensite at -198OC is, by volume fraction,
2% or less of the total residual austenite phase.
[0019] (3) The high strength steel sheet which is
excellent in shapeability according to (1) or (2),
characterized in that the residual austenite phase has a
martensite transformation point of -198OC or less.
[0020] (4) The high strength steel sheet which is
excellent in shapeability according to any one of claims
1 to 3, characterized in that the structure of the steel
sheet further contains, by volume fraction, ferrite
phase: 10 to 75%, bainitic ferrite phase and/or bainite
phase: 10 to 50%, tempered martensite phase: 10 to 50%,
and fresh martensite phase: 10% or less.
[0021] (5) High strength galvanized steel sheet which
is excellent in shapeability characterized by comprising
the high strength steel sheet according to any one of (1)
to (4) on the surface of which a galvanized layer is
formed.
[0022] (6) A method of production of high strength
steel sheet which is excellent in shapeability
characterized by comprising a hot rolling process of
heating a slab which contains, by mass%, C: 0.075 to
0.300%, Si: 0.70 to 2.50%, Mn: 1.30 to 3.50%, P: 0.001 to
0.033%, S: 0.0001 to 0.0100%, Al: 0.005 to 1.500%, N:
0.0001 to 0.0100%, and 0: 0.0001 to 0.0100%, which
contains, as optional elements, one or more of Ti: 0.005
to 0.150%, Nb: 0.005 to 0.150%, B: 0.0001 to 0.0100%, Cr:
0.01 to 2.00%, Ni: 0.01 to 2.00%, Cu: 0.01 to 2.00%, Mo:
0.01 to 1.00%, V: 0.005 to 0.150%, and one or more of Ca,
Ce, Mg, Zr, Hf, and REM: total 0.0001 to 0.5000%, and has
a balance of iron and unavoidable impurities, directly,
or after cooling once, to 1050°C or more, finishing the
rolling at the Ar3 point or more to obtain a steel sheet,
and coiling it at 500 to 750°C in temperature, a cold
rolling process of pickling the coiled steel sheet, then
cold rolling it by a screwdown rate of a screwdown rate
35 to 75%, and an annealing process of heating the steel
sheet after the cold rolling process up to a maximum
heating temperature of 740 to 1000°C, then cooling by an
average cooling rate from the maximum heating temperature
to 700°C of 1.0 to 10.O°C/sec and by a 700 to 500°C
average cooling rate of 5.0 to %OO°C/sec, next holding at
350 to 450°C for 30 to 1000 seconds, then cooling down to
room temperature and, while cooling from the maximum
heating temperature to room temperature, reheating from
the Bs point or less than 500°C to 500°C or more at least
once and reheating from the Ms point or less than 350°C to
350°C or more at least once.
[0023] (7) The method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized by producing high strength steel sheet by
the method of production of high strength steel sheet
according to (6), then galvanizing it.
[0024] (8) A method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized by producing high strength steel sheet by
the method of production according to (6) during the
annealing process of which, at the time of cooling from
the maximum heating temperature to room temperature,
dipping the steel sheet after the cold rolling process in
a zinc bath so as to hot dip galvanize it.
[0025] (9) A method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized by producing high strength steel sheet by
the method of production according to (6) after the
annealing process of which performing hot dip
galvanization.
[0026] (10) A method of production of high strength
galvanized steel sheet which is excellent in shapeability
according to (8) or (9) characterized by performing
alloying treatment at 470 to 650°C in temperature after
the hot dip galvanization.
Advantageous Effects of Invention
[0027] According to the present invention, high
strength steel sheet where a 900 MPa or higher tensile
maximum strength is secured while excellent shapeability
is obtained can be realized.
Brief Description of Drawings
[0028] FIG. 1A is a view which shows an example of a
cooling pattern in annealing treatment in the method of
production of the present invention.
FIG. 1B is a view which shows another example of a
cooling pattern in annealing treatment in the method of
production of the present invention
FIG. 2 is a view which explains an embodiment of the
present invention and a view which shows the relationship
between a tensile strength TS and a total elongation EL.
FIG. 3 is a view which explains an embodiment of the
present invention and a graph which shows the
relationship between a tensile strength TS and a hole
expansion rate A.
Description of Embodiments
[0029] Below, high strength steel sheet and high
strength galvanized steel sheet which are excellent in
shapeability and methods of production of the same of
embodiments of the present invention will be explained.
Note that the following embodiments are explained in
detail for enabling the gist of the present invention to
be understood better, so unless otherwise indicated, do
not limit the present invention.
[0030] Note that, in the following explanation, the
start temperature at which austenite (y-iron) transforms
to martensite in the process of the drop in temperature
in the production of steel sheet will be referred to as
the "Ms point" while the start temperature at which the
residual austenite in the structure of the high strength
steel sheet of the present invention which is produced
transforms to martensite will be referred to as the "Ms,
point".
[0031] First, the structure of the high strength steel
sheet of the present invention will be explained.
[0032] The steel sheet structure of the high strength
steel sheet of the present invention has a 2 to 20%
residual austenite phase. The residual austenite phase
has an Ms, point of -60°C or less. The residual austenite
phase which is contained in such a steel sheet structure
of the high strength steel sheet of the present invention
is stable even with respect to a plurality of deep
cooling treatments.
[0033] The structure other than the residual austenite
phase is not particularly limited so long as a tensile
maximum strength of 900 MPa or higher in strength can be
secured, but preferably has, by volume fraction in the
steel sheet structure, a ferrite phase: 10 to 75%,
bainitic ferrite phase and/or bainite phase: 10 to 50%,
tempered martensite phase: 10 to 50%, and fresh
martensite phase: 10% or less. By having such a steel
sheet structure, the result becomes high strength steel
sheet which has a more excellent shapeability.
[0034] The phases which can be obtained in the
structure of the steel sheet will be explained below:
[0035] Residual Austenite Phase
The residual austenite phase has the property of greatly
improving the strength and ductility, but in general
forming starting points of fracture and greatly degrading
the stretch flangeability.
[0036] In the structure of the present steel sheet, by
reheating two times as explained later, the defects which
were present in the residual austenite phase and were
liable to form starting sites for martensite
transformation are already consumed and only the
austenite phase with its high degree of cleanliness
selectively remains. AS a result, an extremely stable
residual austenite phase is obtained. Such a residual
austenite phase gradually transforms to martensite along
with deformation, so has the property of not easily
forming starting points of fracture and causing extremely
little deterioration of the stretch flangeability.
[0037] As an indicator of the above-mentioned
stability, the martensite transformation start
temperature (Ms, point) of the residual austenite phase
may be mentioned. Stable residual austenite in which an
austenite phase with high degree of cleanliness remains
does not change in amount of residual austenite even with
dipping in liquid nitrogen for 1 hour, that is, applying
so-called deep cooling treatment. The Ms, point is the
liquid nitrogen temperature (-198OC) or less and is
extremely stable. Furthermore, in general, by repeatedly
apply deep cooling treatment, the residual austenite is
gradually decreased, but in the high strength steel sheet
according to the present invention, the residual
austenite does not decrease and is extremely stable even
if treated for deep cooling five times.
[0038] The steel sheet of the present invention gives
high strength steel sheet with a strength and ductility
which are greatly improved and with a stretch
flangeability which is extremely small in deterioration
by a residual austenite phase with an Ms, point of -60°C
or less present in a volume fraction of 2% or more.
[0039] From the viewpoint of the strength and
ductility, the volume fraction of the residual austenite
phase in the steel sheet structure is preferably 4% or
more, more preferably 6% or more. On the other hand, to
make the volume fraction of the residual austenite phase
in the steel sheet structure over 20%, it is necessary to
add elements such as C or Mn in over the suitable
quantity resulting in the weldability being impaired, so
the upper limit of the residual austenite phase is made
20%.
[0040] In the present invention, the ratio of the
residual austenite phase which transforms to martensite
at -198OC is preferably a volume fraction of 2% or less.
Due to this, a more stable residual austenite phase is
obtained, so the ductility and stretch flangeability are
remarkably improved and excellent shapeability is
obtained.
[0041] Further, if the Ms, point of the residual
austenite in the steel sheet structure is -198°C or less,
the result becomes a more stable residual austenite
phase, the ductility and stretch flangeability are
further remarkably improved, and excellent shapeability
is obtained, so this is preferable.
[0042] The volume fraction of the residual austenite
phase is obtained by examining the steel sheet at the
plane parallel to the sheet surface at 1/4 thickness by
X-ray analysis, calculating the area fraction, and
deeming this as the volume fraction. However, the 1/4
thickness plane is made the plane obtained by grinding
and chemically polishing the base material again after
deep cooling treatment to obtain a mirror finish.
[0043] Further, considering measurement error, the
residual austenite phase is deemed to transform to
martensite at the point of time when the relationship
shown below is satisfied:
[0044] Vy(n)/Vy(O) < 0.90
where, "n" is the number of times of deep cooling
treatment, Vy(n) is the residual austenite percent after
the n-th deep drawing treatment, and Vy(0) is the residual
austenite percent in the base material.
[0045] Ferrite Phase
The ferrite phase is a structure which is effective for
improving the ductility and is preferably contained in
the steel sheet structure in a volume fraction of 10 to
75%. If the volume fraction of the ferrite phase in the
steel sheet structure is less than lo%, sufficient
ductility is liable to not be obtained. The volume
fraction of the ferrite phase in the steel sheet
structure, from the viewpoint of the ductility, is more
preferably 15% or more, still more preferably 20% or
more. The ferrite phase is a soft structure, so if the
volume fraction exceeds 75%, sufficient strength will
sometimes not be obtained. To sufficiently raise the
tensile strength of steel sheet, the volume fraction of
the ferrite phase in the steel sheet structure is more
preferably made 65% or less, still more preferably made
50% or less.
[0046] Bainitic Ferrite Phase and/or Bainite Phase
The bainitic ferrite phase and/or bainite phase is a
structure with a good balance of strength and ductility
and is preferably contained in the steel sheet structure
in a volume fraction of 10 to 50%. The bainitic ferrite
phase and/or bainite is a microstructure which has a
strength intermediate to that of a soft ferrite phase and
hard martensite phase and tempered martensite phase and
residual austenite phase. From the viewpoint of the
stretch flangeability, inclusion of 15% or more is more
preferable and inclusion of 20% or more is further
25 preferable. If the volume fraction of the bainitic
ferrite phase and/or bainite exceeds 50%, the yield
stress will excessively rise and the shape freezability
will deteriorate, so this is not preferred.
[0047] Tempered Martensite Phase
30 The tempered martensite phase is a structure which
greatly improves the tensile strength and may be included
in the steel sheet structure to a volume fraction of 50%
or less. From the viewpoint of the tensile strength, the
volume fraction of the tempered martensite is preferably
35 10% or more. If the volume fraction of the tempered
martensite which is contained in the steel sheet
structure exceeds 50%, the yield stress will excessively
rise and the shape freezability deteriorates, so this is
not preferable.
[0048] Fresh Martensite Phase
The fresh martensite phase has the effect of greatly
improving the tensile strength. However, it forms
starting points of fracture and greatly degrades the
stretch flangeability, so it preferably limited to a
volume fraction of 15% in the steel sheet structure. To
raise the stretch flangeability, it is more preferable to
make the volume fraction of the fresh martensite phase in
the steel sheet structure 10% or less, still more
preferably 5% or less.
[0049] Others
The steel sheet structure of the high strength steel
sheet of the present invention may further contain a
pearlite phase and/or coarse cementite phase or other
structure. However, if the steel sheet structure of high
strength steel sheet contains a large amount of pearlite
phase and/or coarse cementite phase, the problem arises
of the bendability deteriorating. From this, the volume
fraction of the pearlite phase and/or coarse cementite
phase which is contained in the steel sheet structure is
preferably a total of 10% or less, more preferably 5% or
less.
[OOSO] The volume fractions of the different
structures which are contained in the steel sheet
structure of the high strength steel sheet of the present
invention can, for example, be measured by the following
method:
[0051] In measuring the volume fractions of the
ferrite phase, bainitic ferrite phase, bainite phase,
tempered martensite phase, and fresh martensite phase
which are contained in the steel sheet structure of the
high strength steel sheet of the present invention,
first, a sample is taken using the cross-section of sheet
thickness parallel to the rolling direction of the steel
sheet as the examined surface. Further, the examined
surface of this sample is polished and etched by Nital
and the range from 1/8 to 3/8 of the sheet thickness is
observed by a field emission scanning electron microscope
(FE-SEM) to measure the area fraction. This was deemed as
the volume fraction.
[0052] Next, the composition of ingredients of the
high strength steel sheet of the present invention will
be explained. Note that in the following explanation
unless particularly designated otherwise, " % " indicates
"mass%".
[0053] C: 0.075 to 0.300%
C is an element which is required for obtaining a
residual austenite phase. It is included for achieving
both an excellent shapeability and high strength. If the
content of C exceeds 0.300%, the weldability becomes
insufficient. From the viewpoint of the weldability, the
content of C is more preferably 0.250% or less, still
more preferably 0.220% or less. If the content of C is
less than 0.075%, it becomes difficult to obtain a
sufficient amount of residual austenite phase and the
strength and shapeability fall. From the viewpoint of the
strength and shapeability, the content of C is more
preferably 0.090% or more, still more preferably 0.100%
or more.
[0054] Si: 0.70 to 2.50%
Si is an element which enables the residual austenite
phase to be easily obtained by suppressing the formation
of iron-based carbides in the steel sheet and is an
element which is necessary for raising the strength and
shapeability. If the content of Si exceeds 2.50%, the
steel sheet becomes brittle and the ductility
deteriorates. From the viewpoint of the ductility, the
content of Si is more preferably 2.20% or less, still
more preferably 2.00% or less. If the content of Si is
less than 0.70%, iron-based carbides form after annealing
while cooling down to room temperature, the residual
austenite phase cannot sufficiently be obtained, and the
strength and shapeability deteriorate. From the viewpoint
of the strength and shapeability, the lower limit value
of Si is more preferably 0.90% or more, still more
preferably 1.00% or more.
[0055] Mn: 1.30 to 3.50%
Mn is added for raising the strength of the steel sheet.
If the content of Mn exceeds 3.50%, coarse MN
concentrated parts form at the center of thickness of the
steel sheet, embrittlement easily occurs, and the
cracking of the cast slab or other trouble easily arises.
Further, if the content of Mn exceeds 3.50%, there is the
problem that the weldability also deteriorates.
Therefore, the content of Mn has to be made 3.50% or
less. From the viewpoint of the weldability, the content
of Mn is more preferably 3.20% or less, still more
preferably 3.00% or less. If the content of Mn is less
than 1.30%, a large amount of soft structures are formed
during the cooling after the annealing, so securing a 900
MPa or more tensile maximum strength becomes difficult.
Therefore, the content of Mn has to be made 1.30% or
more. Further, to raise the strength of the steel sheet,
the content of Mn is more preferably 1.50% or more, still
more preferably 1.70% or more.
[0056] P: 0.001 to 0.030%
P tends to segregate at the center of thickness of the
steel sheet and has the probability of causing the weld
zone to become brittle. If the content of P exceeds
0.030%, the weld zone becomes greatly brittle, so the
content of P is limited to 0.030% or less. The lower
limit of P is not particularly set so long as the effect
of the present invention is exhibited, but if making the
content of P less than 0.001%, the manufacturing cost
greatly increases, so 0.001% is made the lower limit.
[0057] S: 0.0001 to 0.0100%
S has a detrimental effect on the weldability and the
manufacturability at the time of casting and at the time
of hot rolling. Therefore, the upper limit value of the
content of S is made 0.0100% or less. Further, S bonds
with Mn to form coarse MnS which causes the ductility and
stretch flangeability to fall, so the content is more
preferably made 0.0050% or less, still more preferably
0.0025% or less. The lower limit of the content of S is
not particularly set so long as the effect of the present
invention is exhibited, but if making the content of S
less than 0.0001%, the manufacturing cost greatly
increases, so 0.0001% is made the lower limit.
[0058] Al: 0.005 to 1.500%
A1 is an element which suppresses the formation of ironbased
carbides and enables residual austenite to be
easily obtained. It raises the strength and shapeability
of steel sheet. If the content of A1 exceeds 1.500%, the
weldability deteriorates, so the upper limit is made
1.500%. From the viewpoint of the weldability, the
content of A1 is more preferably 1.200% or less, still
more preferably 0.900% or less. A1 is an element which is
effective also as a deoxidizing material, but if the
content of A1 is less than 0.005%, the effect as a
deoxidizing material is not sufficiently obtained, so the
lower limit of the content of A1 is made 0.005%. To
sufficiently obtain the effect of deoxidation, the amount
of A1 is more preferably made 0.010% or more.
[0059] N: 0.0001 to 0.0100%
N forms coarse nitrides which cause the ductility and
stretch flangeability to deteriorate, so the amount of
addition has to be kept down. If the content of Ni
exceeds 0.0100%, this tendency becomes more marked, so
the upper limit of the content of N is made 0.0100%. N
becomes a cause of formation of blowholes at the time of
welding, so the smaller the content, the better. The
lower limit of the content of N is not particularly set
so long as the effect of the present invention is
exhibited, but if making the content of N less than
0.0001%, the manufacturing cost greatly increases, so
0.0001% is made the lower limit.
[0060] 0: 0.0001 to 0.0100%
0 forms oxides which cause the ductility and stretch
flangeability to deteriorate, so the content has to be
kept down. If the content of 0 exceeds 0.0100%, the
deterioration of the stretch flangeability becomes
remarkable, so the upper limit of the content of 0 is
made 0.0100% or less. The content of 0 is more preferably
0.0080% or less, still more preferably 0.0060% or less.
The lower limit of the content of 0 is not particularly
set so long as the effect of the present invention is
exhibited, but if aaking the content of 0 less than
0.0091%, the manufacturing cost greatly increases, so
0.0001% is made the lower limit.
[0061] The high strength steel sheet of the present
invention may further contain the elements which are
shown below in accordance with need:
[0062] Ti: 0.005 to 0.150%
Ti is an element which contributes to the rise in
strength of the steel sheet through precipitation
strengthening, fine grain strengthening by suppression of
growth of ferrite crystal grains, and dislocation
strengthening through suppression of recrystallization.
If the content of Ti exceeds 0.150%, precipitation of
carbonitrides increases and the shapeability
deteriorates, so the content of Ti is made 0.150% or
less. From the viewpoint of the shapeability, the content
of Ti is more preferably 0.100% or less, still more
preferably 0.070% or less. To sufficiently obtain the
effect of the rise in strength by Ti, the content of Ti
has to be made 0.005% or more. To raise the strength of
the steel sheet, the content of Ti is preferably 0.010%
or more, more preferably 0.015% or more.
[0063] Nb: 0.005 to 0.150%
Nb is an element which contributes to the rise in
strength of the steel sheet through precipitation
strengthening, fine grain strengthening by suppression of
growth of ferrite crystal grains, and dislocation
strengthening through suppression of recrystallization.
If the content of Nb exceeds 0.150%, precipitation of
carbonitrides increases and the shapeability
deteriorates, so the content of Nb is made 0.150% or
less. From the viewpoint of the shapeability, the content
of Nb is more preferably 0.100% or less, still more
preferably 0.060% or less. To sufficiently obtain the
effect of the rise in strength by Nb, the content of Nb
has to be made 0.005% or more. To raise the strength of
the steel sheet, the content of Nb is preferably 0.010%
or more, more preferably 0.015% or more.
[0064] V: 0.005 to 0.150%
V is an element which contributes to the rise in strength
of the steel sheet by precipitation strengthening, fine
grain strengthening by suppression of growth of ferrite
crystal grains, and dislocation strengthening through
suppression of recrystallization. If the content of V
exceeds 0.150%, precipitation of carbonitrides increases
and the shapeability deteriorates, so the content is made
0.150% or less. To sufficiently obtain the effect of
raising the strength by V, the content has to be 0.005%
or more.
[0065] B: 0.0001 to 0.0100%
B is an element which suppresses phase transformation at
a high temperature and is effective for increasing the
strength and can be added in place of part of the C
and/or Mn. If the content of B exceeds 0.0100%, the
workability while hot is impaired and the productivity
falls, so the content of B is made 0.0100% or less. From
the viewpoint of the productivity, the content of B is
preferably 0.0050% or less, more preferably 0.0030% or
less. To sufficiently obtain higher strength by B, the
content of B has to be made 0.0001% or more. To
effectively increase the strength of the steel sheet, the
content of B is preferably 0.0003% or more, more
preferably 0.0005% or more.
[0066] Mo: 0.01 to 1.00%
Mo is an element which suppresses phase transformation at
a high temperature and is effective for increasing the
strength and can be added in place of part of the C
and/or Mn. If the content of Mo exceeds 1.00%, the
workability when hot is impaired and the productivity
falls, so the content of Mo is made 1.00% or less. To
sufficiently obtain higher strength by Mo, the content
has to be 0.01% or more.
[0067] W: 0.01 to 1.00%
W is an element which suppresses phase transformation at
a high temperature and is effective for increasing the
strength and can be added in place of part of the C
and/or Mn. If the content of W exceeds 1.00%, the
workability when hot is impaired and the productivity
falls, so the content of W is made 1.00% or less. To
sufficiently obtain higher strength by W, the content has
to be 0.01% or more.
[0068] Cr: 0.01 to 2.00%
Cr is an element which suppresses phase transformation at
a high temperature and is effective for increasing the
strength and can be added in place of part of the C
and/or Mn. If the content of Cr exceeds 2.00%, the
workability when hot is impaired and the productivity
falls, so the content of Cr is made 2.00% or less. To
sufficiently obtain higher strength by Cr, the content
has to be 0.01% or more.
[0069] Ni: 0.01 to 2.00%
Ni is an element which suppresses phase transformation at
a high temperature and is effective for increasing the
strength and can be added in place of part of the C
and/or Mn. If the content of Ni exceeds 2.00%, the
weldability is impaired, so the content of Ni is made
2.00% or less. To sufficiently obtain higher strength by
Ni, the content has to be 0.01% or more.
[0070] Cu: 0.01 to 2.00%
Cu is an element which raises the strength by presence as
fine particles in the steel and can be added in place of
part of the C and/or Mn. If the content of Cu exceeds
2.00%, the weldability is impaired, so the content is
made 2.00% or less. To sufficiently obtain higher
strength by Cu, the content has to be 0.01% or more.
[0071] One or More of Ca, Ce, Mg, Zr, Hf, and REM:
Total 0.0001 to 0.5000%
Ca, Ce, Mg, Zr, Hf, and REM are elements which are
effective for improving the shapeability. One or more can
be added. If the content of the one or more of Ca, Ce,
Mg, Zr, Hf, and REM exceeds a total of 0.5000%,
conversely the ductility is liable to be impaired, so the
total of the contents of the elements is made 0.5000% or
less. To sufficiently obtain the effect of improvement of
the shapeability of the steel sheet, the total of the
contents of the elements has to be 0.0001% or more. From
the viewpoint of the shapeability, the total of the
contents of the elements is preferably 0.0005% or more,
more preferably 0.0010% or more. Here, "REM" is an
abbreviation for "rare earth metal" and indicates the
elements which belong to the lanthanoid series. In the
present invention, the REM or Ce is often added as a
Misch metal. Sometimes, elements of the lanthanoid series
in addition to La or Ce are contained compositely.
Further, even when elements of the lanthanoid series
other than La and Ce are included, the effects of the
present invention are exhibited. Further, even if adding
metal La or Ce, the effects of the present invention are
exhibited.
[0072] Above, the composition of ingredients of the
present invention was explained, but so long as in a
range not impairing the properties of the steel sheet of
the present invention, for example, elements other than
the essential added elements may also be included as
impurities which are derived from the starting materials.
[0073] The high strength steel sheet of the present
invention can also be made high strength galvanized steel
sheet on the surface of which a galvanized layer or
galvannealed layer is formed. By forming a galvanized
layer on the surface of the high strength steel sheet,
steel sheet which has excellent corrosion resistance
results. Further, by forming a galvannealed layer on the
surface of the high strength steel sheet, steel sheet
which has excellent corrosion resistance and which has
excellent coating adhesion results.
100741 Next, the method of production of the high
strength steel sheet of the present invention will be
explained.
[0075] To produce the high strength steel sheet of the
present invention, first, a slab which has the abovementioned
composition of ingredients is cast. As the slab
which is used for hot rolling, for example, it is
possible to use a continuously cast slab or a slab which
is produced by a thin slab caster etc. For the method of
production of the high strength steel sheet of the
present invention, it is preferable to use a process such
as continuous casting-direct rolling (CC-DR) where the
steel is cast, then immediately hot rolled.
[0076] The slab heating temperature in the hot rolling
process has to be 1050°C or more. If the slab heating
temperature is low, the finish rolling temperature falls
below the Ar3 point. As a result, two-phase rolling of the
ferrite phase and austenite phase results, so the hot
rolled sheet structure becomes an uneven mixed grain
structure. The uneven structure is not eliminated even
after the cold rolling and annealing process and
therefore the ductility and bendability deteriorate.
Further, if the finish rolling temperature falls, the
rolling load increases and the rolling becomes difficult
or shape defects are liable to be invited in the steel
sheet after rolling. The upper limit of the slab heating
temperature is not particularly set so long as the effect
of the present invention is exhibited, but it is not
preferable economically to set the heating temperature to
an excessively high temperature, so the upper limit of
the slab heading temperature is preferably made 1350"~ or
less.
[0077] The Ar3 point can be calculated by the following
formula:
[0078] Ar3 ("C) =901-325xC+33xSi-
92x(Mn+Ni/2+Cr/2+Cu/2+Mo/2)+52xAl
[0079] In the above formula, C, Si, Mn, Ni, Cr, Cu,
Mo, and A1 are the contents of the different elements
(mass%) .
[0080] The finish rolling temperature of the hot
rolling is made the higher of 800°C or the Ar3 point as
the lower limit and 1000°C as the upper limit. If the
finish rolling temperature is less than 800°Cf the rolling
load at the time of finish rolling becomes high, the
rolling becomes difficult, and shape defects are liable
to be invited in the hot rolled steel sheet which is
obtained after rolling. If the finish rolling temperature
is less than the Ar3 point, the hot rolling becomes twophase
region rolling of the ferrite phase and austenite
phase and the hot rolled steel sheet structure will
sometimes become an uneven mixed grain structure.
[0081] The upper limit of the finish rolling
temperature is not particularly set so long as the effect
of the present invention is exhibited, but if the finish
rolling temperature is made excessively high, to secure
that temperature, the slab heating temperature has to be
made excessively high. Therefore, the upper limit
temperature of the finish rolling temperature is
preferably made 1000°C or less.
[0082] The steel sheet after rolling is coiled at 500
to 750°C. If coiling the steel sheet at a temperature
exceeding 750°C, the oxides which are formed on the steel
sheet surface excessively increase in thickness and the
pickling ability deteriorates. To raise the pickling
ability, the coiling temperature is preferably 720°C or
less, more preferably 700°C or less. If the coiling
temperature becomes less than 500°C, the hot rolled steel
sheet becomes excessively high in strength and cold
rolling becomes difficult. From the viewpoint of
lightening the load in cold rolling, the coiling
temperature is preferably made 550°C or more. 600°C or
more is more preferable.
[0083] The thus produced hot rolled steel sheet is
pickled. Due to the pickling, the oxides on the steel
sheet surface can be removed. This is important from the
point of improving the chemical convertability of the
cold rolled high strength steel sheet of the final
product or the hot dip coatability of cold rolled steel
sheet for hot dip galvanized or galvannealed steel sheet
use. The pickling may be just a single treatment or may
be divided into a plurality of treatments.
[0084] The pickled steel sheet may be supplied as is
to the annealing process, but by cold rolling it by a
screwdown rate of 35 to 75%, steel sheet with a high
thickness precision and excellent shape is obtained. If
the screwdown rate is less than 35%, it is difficult to
hold the shape flat and the final product becomes poor in
ductility, so the screwdown rate is made 35% or more. If
the screwdown rate exceeds 75%, the cold rolling load
becomes too great and cold rolling becomes difficult.
From this, the upper limit of the screwdown rate is made
75%. The number of rolling passes and the screwdown rate
for each pass are not particularly prescribed so long as
the effect of the present invention is exhibited.
[0085] Next, the obtained hot rolled steel sheet or
cold rolled steel sheet is subjected to the following
annealing treatment.
[0086] First, the rolled steel sheet is heated to a
maximum heating temperature of 740 to 1000°C in range. If
the maximum heating temperature is less than 740°C, the
amount of the austenite phase becomes insufficient and it
becomes difficult to secure a sufficient amount of hard
structures in the phase transformation during subsequent
cooling. If the maximum heating temperature exceeds
1000°C, the austenite phase becomes coarse in particle
size, transformation does not easily proceed during
cooling, and, in particular, a soft ferrite structure
becomes hard to be sufficiently obtained.
[0087] The heating up to the maximum heating
temperature is preferably performed with a heating rate
from the (maximum heating temperature -20)OC to maximum
heating temperature, that is, in the last 20°C at the time
of heating, of 0.1 to 0.8OC/sec. By making the heating
rate in the 20°C up to the maximum heating temperature
gradual heating in the above range, the effects are
obtained that the rate of advance of reverse
transformation to the austenite phase becomes slower and
the defects in the initial austenite phase become
smaller.
[0088] The holding time at the time of heating to the
maximum heating temperature may be suitably determined in
accordance with the maximum heating temperature etc. and
is not particularly limited, but 10 seconds or more is
preferable, while 40 to 540 seconds is more preferable.
[0089] Next, primary cooling is performed by an
average cooling rate from the maximum heating temperature
to 700°C of 1.0 to 10.O°C/sec. By this primary cooling, it
is possible to make the ferrite transformation and
transformation to bainitic ferrite and/or bainite
suitably proceed while leaving a non-transformed
30 austenite phase until the Ms point and transforming all
or part to martensite.
[0090] If the average cooling rate in the above
cooling temperature range is less than l.O°C/sec, pearlite
transformation proceeds during the cooling whereby the
35 nontransformed austenite phase is reduced and a
sufficient hard structure cannot be obtained. As a
result, sometimes it is not possible to secure a tensile
maximum strength 900 MPa or more strength. If the average
cooling rate exceeds 10.O°C/sec, sometimes a soft ferrite
structure cannot be sufficiently formed.
[0091] The holding time in the ferrite transformation
temperature region from right after heating to when the
steel sheet temperature reaches 700°C is not particularly
limited, but is preferably 20 to 1000 seconds. To make
the soft ferrite phase sufficiently form, it is necessary
to hold the steel for 20 seconds or more in the ferrite
transformation temperature region from right after
annealing to when the steel sheet temperature reaches
700°C, preferably hold it there for 30 seconds or more,
more preferably hold it there for 50 seconds or more. If
the time during which the steel is made to held at the
ferrite transformation temperature region exceeds 1000
seconds, the ferrite transformation excessively proceeds,
the nontransformed austenite is reduced, and a sufficient
hard structure cannot be obtained.
[0092] Furthermore, after the above primary cooling,
secondary cooling is performed by an average cooling rate
from 700 to 500°C of 5.0 to 200°C/sec. Due to this
secondary cooling, the transformation from austenite to
ferrite after annealing reliably proceeds. If cooling by
25 a 1°C/sec to 10.O°C/sec average cooling rate similar to
the primary cooling from a temperature region exceeding
700°C, the ferrite phase is insufficiently formed and the
ductility of the high strength steel sheet cannot be
secured.
30 [0093] In the method of production of the present
invention, the steel sheet which has been treated to cool
in the above two stages is held at 350 to 450°C in
temperature for 30 to 1000 seconds of time. If the
holding temperature at this time is less than 350°C, fine
35 iron-based carbides form and concentration of C at the
austenite phase does not proceed resulting in an unstable
austenite phase. If the holding time exceeds 450°C, the
solid solution limit of C in the austenite phase becomes
lower and C becomes saturated even at a small amount, so
concentration of C does not proceed resulting in an
unstable austenite phase.
[0094] If the holding time is less than 30 second, the
bainite transformation does not sufficiently proceed, the
amount of C (carbon) which is discharged from the bainite
phase to the austenite phase is small, the concentration
of C at the austenite phase becomes insufficient, and an
unstable austenite phase results. If the holding time
exceeds 1000 seconds, coarse iron-based carbides start to
form and the concentration of C in the austenite
conversely falls, so an unstable austenite phase results.
[0095] Furthermore, in the annealing process of the
present invention, as shown in FIG. lA, when cooling from
the maximum heating temperature to room temperature, the
steel is reheated from the Bs point (bainite
transformation start temperature) or less than 500°C to
500°C or more at least once and is reheated from the Ms
point or less than 350°C to 350°C or more at least once.
By performing reheating treatment by such two types of
conditions, it is possible to make the austenite phase
which has internal defects and easily transforms to other
structures in the nontransformed residual austenite
phase, that is, the unstable austenite phase,
preferentially transform and obtain a bainite phase,
bainitic ferrite phase, or tempered martensite phase.
[0096] Note that, for example, as shown in FIG. lB,
even if cooling down to the Ms point or less than 350°C,
then reheating to 500°C or more, it is deemed that the
reheating from the Ms point or less than 350°C to 350°C or
more and the reheating from the Bs point or less than
500°C to 500°C or more have respectively been performed.
Such a pattern of reheating treatment may also be
performed.
[0097] Further, it is possible to hold the steel at
the above-mentioned 350 to 450°C temperature range between
the reheating from the Ms point or less than 350°C to
350°C or more and the reheating from the Bs point or less
than 500°C to 500°C or more.
[0098] The Bs point (bainite transformation start
temperature) can be calculated by the following formula:
Bs ("C) =82O-29OC/ (1-VF) -37Si-9OMn-65Cr-5ONi+7OAl
In the above formula, VF is the volume fraction of
ferrite, while C, Mn, Cr, Ni, Al, and Si are the amounts
of addition of these elements (mass % ) .
[0099] The Ms point (martensite transformation start
temperature) can be calculated by the following formula:
Ms ("C) =541-474C/ (1-VF) -15Si-35Mn-17Cr-l7Ni+19Al
In the above formula, VF is the volume fraction of
ferrite, while C, Si, Mn, Cr, Ni, and A1 are the amounts
of addition of these elements (mass % ) .
[OlOO] Note that, it is difficult to directly measure
the volume fraction of the ferrite phase during
pro&~ction of high strength steel sheet, so in the
present invention, a small piece of the cold rolled steel
sheet is cut out before running the sheet through the
continuous annealing line, that small piece is annealed
by the same temperature history as the case of running it
through the continuous annealing line, the change in
volume of the ferrite phase of the small piece is
measured, the result is used to calculate a numerical
value, and that value is used as the volume fraction VF
of the ferrite. This measurement may be performed using
the result of the first measurement operation when
producing steel sheet under the same conditions. The
value does not have to be measured each time. Measurement
is performed again when greatly changing the production
conditions. Of course, it is also possible to observe the
microstructure of the actually produced steel sheet and
feed back the results to the production the next time and
on.
[OlOl] In the above-mentioned reheating from the Bs
point or less than 500°C to 500°C or more, the start
temperature is made the Bs point or less than 500°C so as
to cause the formation of bainite nuclei for consuming
the defects in the austenite. The reheating temperature
was made 500°C or more so as to deactivate the
transformation nuclei and avoid the formation of ironbased
carbides induced by transformation excessively
proceeding in the high temperature region.
[0102] In the above-mentioned reheating from the Ms
point or less than 350°C to 350°C or more, the start
temperature is made the Ms point or less than 350°C so as
to cause the formation of martensite nuclei for consuming
the defects in the austenite. The reheating temperature
was made 350°C or more so as to avoid the formation of
fine iron--based carbides obstructing the concentration of
C at the austenite phase in the martensite and/or bainite
due to being allowed to stand at less than 350°C.
[0103] The reason why performing the above-mentioned
two-stage reheating in different temperature regions
results in the residual austenite phase strikingly rising
is not fully clear, but it is believed that the bainite
nuclei and the martensite nuclei consume respectively
different types of defects.
[0104] Due to the above processes, defects which can
form starting sites of martensite transformation which
are present in the residual austenite phase are consumed,
only the austenite phase with its high degree of
cleanliness selectively remains, and an extremely stable
residual austenite phase is obtained. As a result, high
strength steel sheet which has high ductility and stretch
flangeability and which is excellent in shapeability is
obtained.
[0105] The annealed steel sheet may be cold rolled by
about 0.03 to 0.80% for the purpose of correcting the
shape. At that time, if the cold rolling rate after
annealing is too high, the soft ferrite phase will be
work hardened and the ductility will greatly deteriorate,
so the rolling rate is preferably made the above range.
[0106] The annealed steel sheet may be
electrolytically galvanized to obtain high strength
galvanized steel sheet. Further, the annealed steel sheet
may be hot dip galvanized to obtain high strength
galvanized steel sheet. In such a case, for example, it
is possible to cool from the maximum heating temperature
to room temperature in the annealing process, for
example, down to 500°C, apply further reheating, then dip
in a zinc bath for hot dip galvanization.
[0107] Further, during the secondary cooling during
the above annealing treatment and while holding between
350 to 450°C or after holding at 350 to 450°C, the steel
sheet may be dipped in a galvanization bath to produce
high strength galvanized steel sheet.
[0108] After the hot dip galvanization, it is possible
to further treat the plating layer of the steel sheet
surface to alloy it at a temperature of 470 to 650°C. By
performing such alloying treatment, a Zn-Fe alloy
obtained by the galvanized layer being alloyed is formed
on the surface, and high strength galvanized steel sheet
which is excellent in rust prevention is obtained.
[0109] This heating in the alloying treatment may be
performed in place of the reheating from the Bs point or
less than 500°C to 500°C or more or the reheating from the
Ms point or less than 350°C to 350°C or more.
[OllO] In performing the plating treatment, to improve
the plating adhesion, for example, it is possible to
plate the steel sheet before the annealing process by
plating comprised of one or more elements selected from
Ni, Cu, Co, and Fe. By performing such plating treatment
by this method, high strength galvanized steel sheet
which is formed with a galvanized layer on its surface,
has high ductility and stretch flangeability, and has
excellent shapeability is obtained.
[Olll] The high strength steel sheet on the surface of
which a galvanized layer is formed may further be formed
with a film comprised of a P oxide and/or P-containing
composite oxide.
Examples
[0112] Below, the high strength steel sheet and high
strength galvanized steel sheet which are excellent in
shapeability and the methods of production of the same
of the present invention will be explained more
specifically using examples. The present invention is not
of course limited to the following examples and may be
suitably changed in a range able to match with the gist
of the present invention. These are all included in the
technical scope of the present invention.
[0113] Slabs which have the chemical ingredients
(compositions) of A to AG which are shown in Tables 1 and
2 were cast, then immediately after casting were hot
rolled, cooled, coiled, and pickled under the conditions
which are shown in Tables 3 to 5. After that, Experiments
5, 14, 19, 24, 29, 34, 39, 44, 49, 54, 59, 98, 102, and
119 left the hot rolled steel sheets as they were, while
the other experiments cold rolled them under the
conditions which are described in Tables 3 to 6 after
pickling. After that, an annealing process was applied
under the conditions which are shown in Tables 7 to 14 to
obtain the steel sheets of Experiments 1 to 127.
[0114] Table 1
0
mass%
0.0006
0.0016
0.0025
0.0029
0.0011
0.0022
0.0015
0.0019
0.0020
0.0025
0.0004
0.0013
0.0016
0.0016
0.0004
0.0006
0.0028
0.0016
0.0011
0.0011
0.0019
0.0023
0.0007
0.0007
0.0013
0.0029
0.0014
0.0008
0.0014
0.0009
0.0005
0.0014
0.0008
Experiment
A
B
C
D
E
F
G
H
I
J
K
L
M
N
0
P
Q
R
S
T
U
V
W
X
Y
Z
AA
AB
AC
AD
AE
AF
AG
C
mass%
0.107
0.193
0.107
0.247
0.191
0.133
0.203
0.182
0.084
0.260
0.199
0.094
0.183
0.170
0.140
0.099
0.230
0.119
0.225
0.142
0.194
0.133
0.090
0.101
0.114
0.150
0.015
0.097
0.101
0.093
0.152
0.148
0.209
S i
mass%
1.33
1.97
0.99
1.14
1.05
1.89
1.02
0.75
1.51
0.71
1.19
0.90
2.00
1.66
0.74
0.98
1.24
1.39
1.80
0.99
1.24
2.27
1.44
1.95
1.62
1.06
1.05
0.06
1.05
1.68
0.75
1.72
0.89
A1
mass%
0.043
0.027
0.038
0.005
0.067
0.038
0.073
0.263
0.123
0.302
0.041
0.056
0.045
0.015
0.598
0.044
0.068
0.031
0.032
0.068
0.053
0.071
0.054
0.062
0.071
0.055
0.027
0.027
0.033
0.033
0.065
0.059
0.039
N
mass%
0.0035
0.0021
0.0030
0.0050
0.0030
0.0041
0.0024
0.0037
0.0013
0.0040
0.0060
0.0051
0.0041
0.0037
0.0055
0.0034
0.0054
0.0060
0.0029
0.0021
0.0044
0.0056
0.0020
0.0058
0.0020
0.0018
0.0035
0.0032
0.0033
0.0076
0.0015
0.0080
0.0057
Mn
mass%
1.56
2.49
2.02
1.92
1.41
1.92
1.51
1.87
2.79
2.20
1.89
1.85
1.99
2.59
1.45
1.89
1.45
2.27
1.52
2.17
1.45
2.55
1.68
1.54
2.70
3.16
2.00
1.97
0.52
2.67
2.07
1.55
2.50
P
mass%
0.020
0.014
0.017
0.019
0.015
0.010
0.014
0.012
0.018
0.019
0.018
0.014
0.007
0.020
0.013
0.020
0.016
0.016
0.014
0.011
0.011
0.017
0.016
0.009
0.010
0.010
0.013
0.012
0.015
0.002
0.013
0.007
0.007
S
mass%
0.0038
0.0009
0.0024
0.0034
0.0029
0.0046
0.0052
0.0037
0.0031
0.0014
0.0027
0.0034
0.0018
0.0008
0.0043
0.0007
0.0010
0.0019
0.0042
0.0046
0.0015
0.0051
0.0044
0.0025
0.0034
0.0036
0.0022
0.0022
0.0021
0.0013
0.0018
0.0025
0.0036

[0116] Table 3
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Coiling
temp.
OC
584
600
612
583
5 62
618
638
567
528
632
67 1
695
614
605
660
546
556
613
568
598
682
563
645
62 0
652
639
542
687
577
681
619
661
594
594
643
Experi
-ment
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
2 0
2 1
2 2
2 3
2 4
2 5
2 6
2 7
2 8
2 9
3 0
3 1
3 2
3 3
3 4
3 5
Cold
rolling
rate
9-
0
5 2
5 2
5 2
5 2
0
4 0
4 0
4 0
4 0
52
52
5 2
5 2
0
52
5 2
5 2
52
0
3 8
3 8
67
67
0
5 0
5 0
5 0
50
0
5 2
52
52
52
0
3 8
'lab
heating
temp-
OC
1265
1215
1185
1265
1225
1195
1170
1240
1185
1205
1200
1175
1205
1245
1190
1275
1235
1250
1195
1225
1240
1240
1245
1270
1180
1230
1215
1215
1210
1170
1180
1200
1230
1255
1235
Chemical
ingredients
A
A
A
A
A
B
B
B
B
C
C
C
C
C
D
D
D
D
D
E
E
E
E
E
F
F
F
F
F
G
G
G
G
G
H
Ar 3
transformation
point
OC
769
7 69
769
769
769
67 6
67 6
67 6
67 6
715
715
715
715
715
682
682
682
682
682
747
747
747
747
747
7 4 5
7 4 5
7 4 5
7 4 5
7 4 5
734
734
734
734
7 3 4
686
Hot
rolling
end
temp.
OC
915
901
952
952
92 6
943
910
925
92 9
912
900
8 92
885
923
935
904
930
949
905
913
908
898
908
8 92
944
8 93
92 8
894
943
939
933
893
917
931
890
[0117] Table 4
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp.ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Coiling
temp.
OC
554
572
557
570
528
643
5 5 9
566
615
612
618
638
677
588
658
615
672
611
597
616
669
693
67 9
582
660
5 8 1
569
605
675
5 8 0
628
68 1
613
690
67 1
Experiment
3 6
37
3 8
3 9
4 0
4 1
4 2
4 3
4 4
4 5
4 6
4 7
4 8
4 9
5 0
5 1
52
5 3
5 4
5 5
5 6
5 7
5 8
5 9
60
61
6 2
63
64
65
6 6
6 7
68
6 9
7 0
Cold
rolling
rate
%
3 8
38
3 8
0
68
68
68
6 8
0
3 6
3 6
36
52
0
7 1
7 1
71
7 1
0
5 0
5 0
50
5 0
0
52
52
52
52
52
52
52
52
5 2
52
52
Chemical
ingredients
H
H
H
H
I
I
I
I
I
J
J
J
J
J
K
K
K
K
K
L
L
L
L
L
M
M
M
M
N
N
N
N
0
0
0
Hot
rolling
end
temp.
OC
886
942
92 9
905
8 8 5
951
92 6
953
910
949
938
898
856
933
930
930
942
914
950
9 19
950
902
891
92 4
944
931
8 8 3
945
933
8 95
925
9 14
949
877
903
'lab
heating
temp-
OC
1240
1225
1245
1215
1205
1175
1205
1265
1235
1265
1215
1250
1295
1215
1205
1230
1195
1265
1240
1190
1190
1270
1200
1230
1270
1180
1255
1245
1185
1225
1265
1220
1225
1255
1220
Ar 3
transformation
point
OC
68 6
686
686
686
673
673
673
673
673
642
642
642
642
642
7 0 4
7 0 4
704
704
704
7 10
7 10
7 10
7 10
710
697
697
697
697
652
652
652
652
704
704
704
[0118] Table 5
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Cornp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Cold
rolling
rate
%
52
52
52
52
5 2
52
52
52
52
4 0
4 0
52
52
52
52
5 2
5 2
4 7
4 7
4 7
4 7
67
67
5 2
5 2
52
52
0
52
52
52
0
5 2
52
52
Coiling
temp.
OC
684
609
603
639
614
64 6
67 1
664
591
642
62 9
568
602
648
573
578
589
625
652
585
558
551
608
591
607
67 8
645
650
582
664
657
564
606
634
604
Hot
rolling
end
temp .
OC
915
932
923
903
927
890
934
913
947
9 0 9
907
886
924
900
9 18
9 3 1
920
942
889
907
8 97
904
904
897
930
909
8 9 9
901
917
8 8 8
907
92 1
9 14
910
952
Experiment
7 1
7 2
7 3
7 4
7 5
7 6
7 7
7 8
7 9
8 0
8 1
8 2
8 3
8 4
8 5
8 6
8 7
8 8
8 9
9 0
9 1
92
9 3
94
95
9 6
97
9 8
9 9
1CO
10 1
102
10 3
10 4
10 5
'lab
heating
temp-
OC
1215
1230
1180
1230
1215
1180
1270
1260
1280
1190
1245
1205
1210
1215
1180
1210
1265
1245
1275
1275
1230
1225
1190
1205
1275
1185
1200
1215
1230
1260
1190
1195
1280
1235
1275
Chemical
ingredients
0
P
P
P
P
Q
Q
Q
Q
R
R
R
R
S
S
S
S
T
T
T
T
U
U
U
U
V
V
V
V
W
W
W
W
X
X
Ar 3
transformation
point
OC
704
730
730
730
7 3 0
737
737
737
737
7 0 1
701
701
701
7 4 9
749
749
749
691
691
691
691
748
7 4 8
7 4 8
7 4 8
691
691
691
691
768
7 68
7 68
7 68
755
755
[0119] Table 6
Cold
rolling
rate
3-
0
52
52
67
67
4 3
4 3
52
52
52
52
52
52
52
0
5 0
5 0
5 0
Coiling
temp.
OC
62 1
616
673
652
673
5 63
643
694
683
666
62 3
622
660
566
5 8 8
600
601
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Comp. ex.
Comp. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
-
Experiment
10 6
107
10 8
10 9
11 0
111
112
11 3
114
115
116
117
11 8
119
120
121
122
'lab
heating
temp-
OC
1210
1280
1200
1185
1175
1185
1225
1185
1220
1275
1190
1205
1175
1210
1200
1225
1230
Chemical
ingredients
X
X
Y
Y
Y
Y
Z
Z
Z
Z
AA
AB
AC
B
AD
AD
AE
Ar 3
transformation
point
OC
755
755
673
673
67 3
673
5 9 9
5 9 9
599
599
7 4 8
692
857
67 6
682
682
68 9
Hot
rolling
end
temp .
OC
900
939
886
925
940
953
92 9
915
902
-735
935
889
894
925
866
903
889
[0120] Table 7
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Experiment
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
2 0
2 1
2 2
2 3
2 4
2 5
2 6
2 7
2 8
2 9
3 0
3 1
3 2
3 3
3 4
3 5
First
cooling
process
Average
cooling
rate
OC/sec
3.3
3.1
3.1
3.4
3.3
1.7
1.6
1.5
2.5
2.5
2.6
3.2
3.4
3.1
6.2
4.5
3.6
4.4
3.7
2.6
3.1
2.5
2.5
2.6
2.5
3.5
2.9
2.7
2.7
3.3
2.8
3.2
2.7
3.5
4.9
Second
cooling
process
Average
cooling
rate
~C/sec
3 4
32
2 8
2 9
3 3
2 7
2 8
3 1
2 7
13
11
11
10
13
10
10
8
10
8
7
12
8
12
10
8
7
4 6
52
5 5
3 1
3 5
32
3 0
3 0
3 3
Chemical
ingredients
A
A
A
A
A
B
B
B
B
C
C
C
C
C
D
D
D
D
D
E
E
E
E
E
F
F
F
F
F
G
G
G
G
G
H
Steel
type
CR
CR
G A
GI
HR-GA
CR
CR
CR
G A
CR
CR
CR
EG
HR
CR
CR
CR
GI
HR
CR
CR
G A
EG
HR-GA
CR
CR
CR
CR
HR
CR
CR
CR
EG
HR-GA
CR
Heating
Heating
rate
~C/sec
0.3
0.3
0.3
0.3
0.3
0.3
0.3
0.3
0.3
0.5
0.5
-15
0.5
0.4
0.7
0.7
-8
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.7
0.5
0.5
0.5
0.5
0.5
0.5
process
Max.
heating
temp.
OC
92 0
812
816
821
812
819
825
826
823
846
836
831
845
8 4 5
793
782
781
7 8 6
784
822
829
823
821
8 16
834
8 9 8
892
-1076
898
793
7 8 9
-7 3 0
783
8 0 0
780
[0121] Table 8
Inv. ex.
Comp.ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp.ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp-ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Second
cooling
process
Average
cooling
rate
OC/sec
3 1
30
3 0
3 3
2 6
3 4
33
3 5
3 0
3 4
102
- 1
3 3
3 1
5 9
65
57
61
63
5 6
61
58
6 6
5 3
5 7
11
59
64
2 7
3 1
29
3 1
8
3 0
32
First
cooling
process
Average
cooling
rate
OC/sec
5.3
-3 1
5.4
2.6
3.1
3.2
-0.2
2.5
3.5
2.5
3.3
3.6
2.4
2.7
3.3
2.7
2.7
3.1
3.1
2.9
2.9
3.5
2.8
3.1
3.0
3.5
3.4
2.7
3.0
3.1
3.5
2.6
3.8
3.3
3.7
Experiment
3 6
37
3 8
3 9
4 0
4 1
4 2
4 3
4 4
4 5
4 6
4 7
4 8
4 9
5 0
5 1
52
5 3
5 4
5 5
5 6
5 7
5 8
5 9
6 0
61
6 2
63
64
65
6 6
67
68
6 9
7 0
Steel
type
CR
CR
G A
HR
CR
CR
CR
GI
HR-GA
CR
CR
CR
GI
HR-GA
CR
CR
CR
GI
HR-GA
CR
CR
CR
GI
HR-GA
CR
CR
CR
G A
CR
CR
CR
GI
CR
CR
CR
Chemical
ingredients
H
H
H
H
I
I
I
I
I
J
J
J
J
J
K
K
K
K
K
L
L
L
L
L
M
M
M
M
N
N
N
N
0
0
0
Heating
Heating
rate
OC/sec
0.5
0.5
0.5
0.5
0.5
0.6
0.5
0.5
0.5
0.4
0.4
0.4
0.4
0.4
0.5
0.5
0.5
0.5
0.5
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.5
0.5
0.5
process
Max.
heating
temp .
OC
8 8 9
886
788
790
817
8 17
818
811
828
8 4 0
835
8 5 6
835
846
8 10
793
804
796
799
821
837
828
837
826
829
824
823
8 15
827
821
810
8 18
953
943
944
[0122] Table 9
Second
cooling
process
Average
cooling
rate
OC/sec
4 1
2 7
3 0
33
8
2 8
3 0
26
2 9
2 7
3 0
30
3 3
2 7
128
3 5
2 8
32
3 4
32
2 9
2 8
3 2
2 6
2 7
4 7
4 9
52
5 4
5 0
5 1
4 7
5 4
4 7
5 3
First
cooling
process
Average
cooling
rate
OC/sec
2.7
7.8
2.6
2.9
2.6
2.8
3.4
2.7
2.8
3.2
3.3
2.6
2.7
1.9
1.8
2.4
2.1
2.1
1.6
2.4
1.9
4.8
5.4
4.7
5.2
5.0
5.1
4.8
5.2
5.1
5.3
4.6
5.1
5.4
4.9
Inv. ex.
Inv. ex.
Inv. ex.
Comp.ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Experiment
7 1
7 2
7 3
7 4
7 5
7 6
7 7
7 8
7 9
8 0
8 1
8 2
8 3
8 4
8 5
8 6
8 7
8 8
8 9
9 0
9 1
92
9 3
9 4
9 5
9 6
97
9 8
9 9
10 0
101
102
10 3
10 4
10 5
Steel
type
G A
CR
CR
CR
G A
CR
CR
CR
G A
CR
CR
CR
GI
CR
CR
G A
GI
CR
CR
G A
EG
CR
CR
G A
E G
CR
CR
HR
G A
CR
CR
HR-GA
GI
CR
CR
Chemical
ingredients
0
P
P
P
P
Q
Q
Q
Q
R
R
R
R
S
S
S
S
T
T
T
T
U
U
U
U
V
V
V
V
W
W
W
W
X
X
Heating
Heating
rate
OC/sec
0.5
0.4
0.3
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.5
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.4
0.2
0.2
0.2
0.2
0.5
0.5
process
Max.
heating
temp .
OC
939
849
838
842
841
794
801
8 0 0
806
8 17
803
8 0 0
807
798
806
8 0 1
804
835
820
826
833
785
7 7 1
7 8 7
7 7 5
8 65
8 8 0
872
8 67
882
796
793
804
852
8 4 7
[0123] Table 10
nv. ex.
[0124] Table 11
[0125] Table 12
Martensite
transformation
start
temp .
(Ms)
Experimen
Reheating
process 1
Coolsy'p
temp.
Reheat
stop
temp.
Alloying
process
Alloying
temp.
Reheating
process 2
Bainite
transformation
start
temp.
(BS)
ing
stop
temp.
Reheat
temp.
Reheating
process 3
Holding
process
Holding
t ime
Cooling
stop
temp.
Reheat
temp.
Reheating
process 4
Cooling
stop
temp.
Reheating
process 5
Reheat
stop
temp.
Cooling
stop
temp.
Reheat
stop
t emp .

LO1271 Table 14
[0128] In the annealing process, first, the steel
sheets were heated to the maximum heating temperatures
described in Tables 7 to 10 by average heating rates
between the (maximum heating temperature -20°C) to maximum
heating temperature of the average heating rates
described in Tables 7 to 10. Next, in the first cooling
process (primary cooling) from the maximum heating
temperature to 700°C, they were cooled by the average
cooling rates described in Tables 7 to 10. Furthermore,
in the second cooling process (secondary cooling) from
700°C to 500°C, they were cooled by the average cooling
rates described in Tables 7 to 10.
[0129] After that, the steel sheets were reheated from
the Bs point or 480°C or less to 500°C or more 1 to 3
times (reheating processes 1, 2, and 4) and, furthermore,
were reheated from the Ms point or 350°C or less to 350°C
or more 1 or 2 times (reheating processes 3 and 5).
[0130] After the reheating process 3, the steel sheets
were held at 300 to 450°C in range for exactly the times
described in Tables 11 to 14 then were treated by the
reheating processes 4 and 5 and cooled down to room
temperature.
[0131] After being cooled down to room temperature, in
Experiments 6 to 49, the steel sheets were cold rolled by
0.15%, in Experiments 60 to 83, the steel sheets were
cold rolled by 0.30%, in Experiment 89, the steel sheet
was cold rolled by 1.50%, in Experiment 93, the steel
sheet was cold rolled by 1.00%, and in Experiments 96 to
118 and 120 to 127, the steel sheets were cold rolled by
0.25%.
[0132] The types of steels in the experiments are
shown in the tables as cold rolled steel sheet (CR), hot
rolled steel sheet (HR), electrolytically galvanized
steel sheet (EG) , hot dip galvanized steel sheet (GI) ,
hot dip galvannealed steel sheet (GA), and hot rolled hot
dip galvannealed steel sheet (HR-GA) (same in tables
shown below) .
[0133] Experiments 13, 23, 33, 91, 95, 107, and 114
are examples in which the steel sheets were electroplated
after the annealing process to obtain galvanized steel
sheets (EG).
[0134] Experiments 4, 18, 43, 83, and 87 are examples
in which after the second cooling process, the steel
sheets are dipped in a galvanization bath until the
holding treatment at 350 to 450°C in range to obtain hot
dipped galvanized steel sheets (GI).
[0135] Experiments 48, 53, 58, 98, and 103 are
examples in which after the holding treatment at 300 to
450°C in range, the steel sheets are dipped in a
galvanization bath, then are cooled to room temperature
to obtain hot dipped galvanized steel sheets (GI).
[0136] Experiments 3, 5, 9, 34, 38, 44, 49, 67, 86,
90, 94, 99, 102, and 110 are examples in which after the
second cooling process, the steel sheets are dipped in a
galvanization bath until holding at 350 to 450°C in range
and then are treated for alloying at the described
temperatures to obtain galvannealed steel sheets (GA) .
[0137] Experiments 22, 24, 54, 59, 63, 71, 75, 79,
121, 123, 125, and 127 are examples in which after the
holding treatment at 300 to 450°C in range, the steel
sheets are dipped in a galvanization bath and treated for
alloying at the described temperatures to obtain hot
dipped galvannealed steel sheets (GA).
[0138] Experiments 9, 63, and 90 are examples in which
the surfaces of the plating layers are given films
comprised of P-based composite oxides.
[0139] Tables 15 to 18 give the results of analysis of
the microstructures of the steel sheets of Experiments 1
to 127. In the microstructure fractions, the amounts of
residual austenite (residual y) were measured by X-ray
diffraction at planes parallel to the sheet thickness at
1/4 thickness. The rest gives the results of measurement
of the fractions of microstructures in the range of 1/8
thickness to 3/8 thickness. Sheet thickness crosssections
parallel to the rolling direction were cut out,
polished to mirror surfaces, etched by Nital, then
5 examined using field emission scanning electron
microscope (FE-SEM) .
[0140] Table 15
[0141] Table 16
[0142] Table 17
[0143] Table 18
[0144] Tables 19 to 22 show the results of measurement
of the residual austenite fractions and the amounts of
solid solution C in the residual austenite after deep
cooling treatment tests. These were measured by X-ray
diffraction at planes parallel to the sheet thickness at
1/4 thickness. The Ms, points were measured by preparing
liquid nitrogen
(-198OC) and ethanol cooled using liquid nitrogen in 20°C
increments from O°C to -lOO°C, holding the steel sheets at
those temperatures for 1 hour, then measuring the
residual austenite fractions and using the maximum
temperatures at which the austenite fractions fall as the
Ms, points of the residual austenite phase.
[0145] Table 19
Inv. ex.
Experiment
i
type
CR
Chemical
ingredients
A
Residual y transformation temp.
Solid
solution
C amount
0
0
0.93
MS, point
OC
< -198
Volume
fraction after
liquid
nitrogen
dipping
1st
G
3rd
G
5th
G
- 52 -
[0146] Table 20
fraction after
- 53 -
[0147] Table 21
fraction after
[0148] Table 22
[0149] In the dipping treatment in liquid nitrogen,
the operation from dipping the steel sheet in liquid
nitrogen for 1 hour, then taking it out and allowing it
to stand in the air until reaching room temperature is
counted as one treatment. The residual austenite
fractions were measured at the points of time of the ends
of the first, third, and 10th treatments. Steel sheets
with residual austenite fractions which did not change
were evaluated as "G (good)" while steel sheets with
residual austenite fractions which decreased were
evaluated as "P (poor) " .
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Inv. ex.
Comp. ex.
Comp. ex.
Comp. ex.
Comp. ex.
Inv. ex.
Inv. ex.
Inv. ex.
type
CR
EG
CR
CR
G A
CR
CR
CR
E G
CR
CR
CR
CR
HR
CR
G A
Experiment
106
107
10 8
109
110
11 1
112
113
11 4
11 5
11 6
117
11 8
11 9
120
12 1
Chemical
ingredients
X
X
Y
Y
Y
Y
Z
Z
Z
Z
AA
AB
AC
B
AD
AD
temp.
Solid
solution
C amount
0
0
0.89
1.07
0.96
0.93
0.99
1.04
1.01
0.91
1.05
0.95
Residual transformation
MS, point
OC
< -198
< -198
< -198
< -198
< -198
< -198
< -198
-------< -19-8 -
< -198
< -198
No
No residual austenite
No residual austenite
Volume
fraction after
liquid
nitrogen
dipping
1st
G
G
G
G
G
G
G
G
G
< -198
< -198
< -198
G
G
G
G
G
G
residual austenite
3rd
P
G
G
G
G
G
G G G
G
G
G
5th
G
G
G
G
G
G
G
G
G
G
G
0.90
0.99
0.89
[0150] Tables 23 to 26 show the evaluation of
properties of the steel sheets of Experiments 1 to 127.
At that time, tensile test pieces based on JIS Z 2201
were taken from the steel sheets of Experiments 1 to 127
and were subjected to tensile tests based on JIS Z 2241
to measure the yield strength (YS), tensile strength
(TS) , and total elongation (EL) .
[0151] FIG. 2 shows the relationship between the
tensile strength (TS) and the total elongation (EL),
while FIG. 3 shows the relationship between the tensile
strength (TS) and the hole expansion rate (h) which
serves as an indicator of the stretch flangeability. The
steel sheets of the present invention satisfy all of
TS2900MPa, TSxEL>17000MPa.%, TSxh224000MPa-%. The steel
sheets of the comparative examples are not steel sheets
which satisfy all of these.
101521 Table 23
Experiment
1
2
Chemical
ingredients
A
A
3 3
- 34
3 5
Steel
type
CR
CR
G
G
H
Material
measurement results
EG
HR-GA
CR
TS
x
EL
MPae%
19494
19076
YS
MPa
943
503
581
619
665
T S
x
h
MPae%
55404
39086
TS
MPa
1026
1004
Inv. ex.
Inv. ex.
1147
1158
1071
EL
%
19
19
%
54
39
22
17
16
50
48
46
25234
19686
17136
57350
55584
49596
Inv. ex.
Inv. ex.
Inv. ex.
- 57 -
[0153] Table 24
- 58 -
[0154] Table 25
Experiment
7 1
7 2
Chemical
ingredients
0
P
10 3
10 4
105
Steel
type
GA
CR
W
X
X
Material
measurement 'results
GI
CR
CR
TS
x
EL
MPae%
20539
20881
YS
MPa
659
1002
391
764
626
T S
x
h
MPae%
45402
41762
TS
MPa
1081
1099
Inv. ex.
Inv. ex.
1013
1104
1112
EL
%
19
19
h
%
42
38
17
16
18
44
47
41
17221
17664
20016
44572
51610
45592
Inv. ex.
Inv. ex.
Inv. ex.
- 59 -
[0155] Table 26
[0156] Experiment 115 is an example in which the end
temperature of the hot rolling is low. The microstructure
5 is stretched in one direction making it uneven, so the
ductility and stretch flangeability are poor.
[0157] Experiments 12, 17, 106, and 111 are examples
in which the heating rate from the (maximum heating
temperature -20°C) in the heating process is large. The
10 residual austenite phase is unstable and the stretch
flangeability is poor.
[0158] Experiment 28 is an example in which the
maximum heating temperature in the annealing process is
high. The soft structure is not sufficiently formed and
15 the ductility is poor.
Experiment
106
107
126 AG CR 635 11521 19 1 39 1 21888 ) 44928 1 Inv. ex.
127 AG GA 701 10551 21 1 46 1 22155 1 48530 I Inv. ex.
Chemical
ingredients
X
X
Steel
type
CR
EG
Material
measurement results TS
x
EL
MPa.%
19312
21580
YS
MPa
707
663
T S
x
h
MPae%
26201
49634
TS
MPa
1136
1079
Inv. ex.
Inv. ex.
EL
%
17
20
h
%
23
46
[0159] Experiment 32 is an example in which the
maximum heating temperature in the annealing process is
low. A large number of coarse iron-based carbides which
form starting points of fracture are included, so the
ductility and stretch flangeability are poor.
[0160] Experiment 37 is an example in which the
average cooling rate in the first cooling process
(primary cooling) is high. Soft structures are not
sufficiently formed, so the ductility and stretch
flangeability are poor.
[0161] Experiment 42 is an example in which the
average cooling rate in the first cooling process
(primary cooling) is low. Coarse iron-based carbides are
formed, and the stretch flangeability is poor
[0162] Experiment 47 is an example in which the
cooling rate in the second cooling process (secondary
cooling) is low. Coarse iron-based carbides are formed,
and the stretch flangeability is poor.
[0163] Experiment 52 is an example where no reheating
treatment is performed. The residual austenite phase is
unstable, and the stretch flangeability is poor.
[0164] Experiments 57, 66, and 82 are examples where
only reheating from the Bs point or 480°C or less to 500°C
or more is performed. The residual austenite phase is
unstable, and the stretch flangeability is poor.
[0165] Experiments 62 and 70 are examples where only
reheating from the Ms point or 350°C or less to 350°C or
more is performed. The residual austenite phase is
unstable, and the stretch flangeability is poor.
[0166] Experiment 74 is an example where the time of
the treatment at 300 to 450°C in range is short. Carbon
does not concentrate at the residual austenite, the
residual austenite phase is unstable, and the stretch
flangeability is poor.
[0167] Next, Experiment 78 is an example where the
holding time at 300 to 450°C in range is long. Iron-based
carbides form, the amount of solid solution C in the
residual austenite falls, the residual austenite phase is
unstable, and the stretch flangeability is poor.
[0168] Next, Experiments 116 to 118 are examples where
5 the composition of ingredients deviated from the
predetermined range. In each case, sufficient properties
could not be obtained.
[0169] From the results of the examples which were
explained above, it is clear that according to the high
10 strength steel sheet and high strength galvanized steel
sheet which are excellent in shapeability and methods of
prcduction cf the same of the present invention, high
strength steel sheet which secures a tensile maximum
strength of 900 MPa or more in high strength while is
15 given excellent ductility and stretch flangeability and
has sufficiently high shapeability is obtained.
Industrial Applicability
[0170] According to the present invention, for
example, in applications such as members which are
20 obtained by shaping steel sheet by press working etc., a
tensile maximum strength of 900 MPa or more of high
strength is secured while excellent ductility and stretch
flangeability are obtained and excellent strength and
shapeability are simultaneously obtained. Due to this,
25 for example, in particular, by applying the present
invention to the field of auto parts etc., in particular
by applying it to the field of automobiles, it is
possible to fully enjoy the merits of improved safety
along with the increased strength of the chassis,
30 improved shapeability at the time of working the members,
etc. The contribution to society is immeasurable.

C
nngcr ,-a* J& NAL
- 62 -
L.
CLAIMS j -
Claim 1. High strength steel sheet which is
excellent in shapeability which contains, by mass%,
C: 0.075 to 0.300%,
Si: 0.70 to 2.50%,
Mn: 1.30 to 3.50%,
P: 0.001 to 0.030%,
S: 0.0001 to 0.0100%,
Al: 0.005 to 1.500%,
N: 0.0001 to 0.0100%, and
0: 0.0001 to 0.0100%,
which contains, as optional elements, one or more of
Ti: 0.005 to 0.150%,
Nb: 0.005 to 0.150%,
B: 0.0001 to 0.0100%,
Cr: 0.01 to 2.00%,
Ni: 0.01 to 2.00%,
Cu: 0.01 to 2.00%,
Mo: 0.01 to 1.00%,
V: 0.005 to 0.150%, and
one or more of Ca, Ce, Mg, Zr, Hf, and REM: total
0.0001 to 0.5000%, and
has a balance of iron and unavoidable impurities,
wherein
the structure of the steel sheet contains, by volume
fraction, 2 to 20% of residual austenite phase, and
said residual austenite phase has a martensite
transformation point of -60°C or less.
Claim 2. The high strength steel sheet which is
excellent in shapeability according to claim 1,
characterized in that a ratio of said residual austenite
phase which transforms to martensite at -198O~ is, by
volume fraction, 2% or less of the total residual
austenite phase.
Claim 3. The high strength steel sheet which is
excellent in shapeability according to claim 1 or 2,
characterized in that said residual austenite phase has a
- 63 -
martensite transformation point of -198OC or less.
Claim 4. The high strength steel sheet which is
excellent in shapeability according to claim 1 or 2.
characterized in that the structure of the steel sheet
further contains, by volume fraction,
ferrite phase: 10 to 75%,
bainitic ferrite phase and/or bainite phase: 10 to
50%,
tempered martsnsite phase: 10 to 50%, and
fresh martensite phase: 10% or less.
Claim 5. High strength galvanized steel sheet which
is excellent in shapeability characterized by comprising
the high strength steel sheet according to claim 1 or 2
on the surface of which a galvanized layer is formed.
15 Claim 6. A method of production of high strength
steel sheet which is excellent in shapeability
characterized by comprising:
a hot rolling process of heating a slab which
contains, by mass%,
Si: 0.70 to 2.50%,
Mn: 1.30 to 3.50%,
P: 0.001 to 0.030%,
S: 0.0001 to 0.0100%,
Al: 0.005 to 1.500%,
N: 0.0001 to 0.0100%, and
0: 0.0001 to 0.0100%,
which contains, as optional elements, one or more of
Ti: 0.005 to 0.150%,
Nb: 0.005 to 0.150%,
B: 0.0001 to 0.0100%,
Cr: 0.01 to 2.00%,
Ni: 0.01 to 2.00%,
Cu: 0.01 to 2.00%,
Mo: 0.01 to 1.00%,
V: 0.005 to 0.150%, and
0.0001 to 0.5000%, and
has a balance of iron and unavoidable impurities,
directly, or after cooling once, to 1050°C or more,
finishing the rolling at the Ar3 point or more to obtain a
steel sheet, and coiling it at 500 to 750°C in
temperature,
a cold rolling process of pickling the coiled steel
sheet, then cold rolling it by a screwdown rate of a
screwdown rate 35 to 75%, and
an annealing process of heating the steel sheet
after the cold rolling process up to a maximum heating
temperature of 740 to 1000°C, then cooling by an average
cooling rate from said maximum heating temperature to
700°C of 1.0 to 10.O°C/sec and by a 700 to 500°C average
cooling rate of 5.0 to 200°C/sec, next holding at 350 to
450°C for 30 to 1000 seconds, then cooling down to room
temperature and, while cooling from said maximum heating
temperature to room temperature, reheating from the Bs
point or less than 500°C to 500°C or more at least once
and reheating from the Ms point or less than 350°C to
350°C or more at least once.
Claim 7. A method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized by producing high strength steel sheet by
the method of production of high strength steel sheet
according to claim 6, then galvanizing it.
Claim 8. A method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized by producing high strength steel sheet by
the method of production according to claim 6 during the
annealing process of which, at the time of cooling from
said maximum heating temperature to room temperature,
dipping the steel sheet after said cold rolling process
in a zinc bath so as to hot dip galvanize it.
Claim 9. A method of production of high strength
galvanized steel sheet which is excellent in shapeability
characterized -by producing high strength steel sheet by
the method of production according to claim 6 after the
annealing process of which performing hot dip
galvanization.
Claim 10. A method of production of high strength
galvanized steel sheet which is excellent in shapeability
according to claim 8 or 9 characterized by performing
alloying treatment at 470 to 650'~ in temperature after
said hot dip galvanization.
Dated this 3" day of Fe bruary, 201 4
- -
[S WATI PAHuJA]
OF RIMFRY dt SAGAR
ATTORNEY FOR mx