Technical Field
[0001] The present invention relates to a high strength steel sheet excellent in tensile
strength, press formability, flatness in the width direction, and working embrittlement
resistance, and to a method for manufacturing the same. The high strength steel sheet
of the present invention may be suitably used as structural members, such as automobile
parts.
Background Art
[0002] Steel sheets for automobiles are being increased in strength in order to reduce CO
2 emissions by weight reduction of vehicles and to enhance crashworthiness by weight
reduction of automobile bodies at the same time, with introduction of new laws and
regulations one after another. To increase the strength of automobile bodies, high
strength steel sheets having a tensile strength (TS) of 1180 MPa or higher grade are
increasingly applied to principal structural parts of automobiles.
[0003] High strength steel sheets used in automobiles require excellent press formability.
For example, high strength steel sheets with high El and excellent hole expansion
ratio λ are suitably applied to automobile frame parts, such as bumpers. From the
point of view of crash safety, excellent working embrittlement resistance is required.
[0004] Furthermore, high strength steel sheets used in automobiles require high flatness.
Patent Literature 1 describes that warpage of a steel sheet causes operational troubles
in forming lines and adversely affects the dimensional accuracy of products. The present
inventors carried out extensive studies and have found that the dimensional accuracy
of products is affected not only by the warpage of steel sheets but also by the flatness
in the width direction that is evaluated as steepness. For example, the steepness
in the width direction is suitably 0.02 or less in order to achieve excellent dimensional
accuracy.
[0005] To meet the above demands, for example, Patent Literature 2 provides a hot-dip galvanized
steel sheet with excellent press formability and low-temperature toughness that has
a tensile strength of 980 MPa or more, and a method for manufacturing the same. While
the steel sheet of Patent Literature 2 is improved in embrittlement at low temperatures,
the technique does not take into consideration the working embrittlement of the steel
sheet or the flatness in the width direction.
Citation List
Patent Literature
Non Patent Literature
Summary of Invention
Technical Problem
[0008] The present invention has been developed in view of the circumstances discussed above.
Objects of the present invention are therefore to provide a high strength steel sheet
having 980 MPa or higher TS and being excellent in press formability, flatness in
the width direction, and working embrittlement resistance; and to provide a method
for manufacturing the same.
Solution to Problem
[0009] The present inventors carried out extensive studies directed to solving the problems
described above and have consequently found the following facts.
- (1) 980 MPa or higher TS and excellent press formability can be realized by limiting
the amount of tempered martensite to 38% or more and less than 90%, the amount of
the total of ferrite and bainitic ferrite to 10% or more and 60% or less, and the
amount of retained austenite to less than 3%.
- (2) The flatness in the width direction can be enhanced by limiting the proportion
of a packet having the largest area in tempered martensite to 70% or less of a prior
austenite grain.
- (3) Excellent working embrittlement resistance can be achieved by limiting the proportion
of a packet having the largest area in tempered martensite to 70% or less of a prior
austenite grain and by limiting the average prior austenite grain size in tempered
martensite to 20 um or less.
[0010] The present invention has been made based on the above findings. Specifically, a
summary of configurations of the present invention is as follows.
- [1] A high strength steel sheet having a chemical composition including, in mass%,
C: 0.030% or more and 0.500% or less, Si: 0.01% or more and 2.50% or less, Mn: 0.10%
or more and 5.00% or less, P: 0.100% or less, S: 0.0200% or less, Al: 1.000% or less,
N: 0.0100% or less, and O: 0.0100% or less, a balance being Fe and incidental impurities,
the high strength steel sheet being such that in a region at 1/4 sheet thickness,
an area fraction of tempered martensite is 38% or more and less than 90%, a volume
fraction of retained austenite is less than 3%, an area fraction of the total of ferrite
and bainitic ferrite is 10% or more and 60% or less, an average grain size of prior
austenite is 20 um or less, and an average of the proportions of packets having the
largest area in prior austenite grains is 70% by area or less of the prior austenite
grain.
- [2] The high strength steel sheet according to [1], wherein the chemical composition
further includes at least one element selected from, in mass%, Ti: 0.200% or less,
Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100%
or less, Cr: 1.00% or less, Mo: 1.00% or less, Co: 0.010% or less, Ni: 1.00% or less,
Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg:
0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf:
0.10% or less, and Bi: 0.200% or less.
- [3] The high strength steel sheet according to [1] or [2], which has a coated layer
on a surface of the steel sheet.
- [4] A method for manufacturing the high strength steel sheet according to [1] or [2],
the method including providing a cold rolled steel sheet produced by subjecting a
steel having the chemical composition described above to hot rolling, pickling, and
cold rolling; heating the steel sheet at an annealing temperature T1 of 700°C or above
and 950°C or below for a holding time t1 at the annealing temperature T1 of 10 seconds
or more and 1000 seconds or less; cooling the steel sheet in such a manner that the
average cooling rate from 750°C to 600°C is less than 20°C/s, the average cooling
rate from (Ms + 50°C) to a quench start temperature T2 is less than 5°C/s wherein
the quench start temperature T2 is (Ms - 80°C) or above and is below Ms where Ms is
martensite start temperature (°C) defined by formula (1), and the average cooling
rate from the quench start temperature T2 to 80°C is 300°C/s or more; and heating
the steel sheet at a tempering temperature T3 of 100°C or above and 400°C or below
for a holding time t3 at the tempering temperature T3 of 10 seconds or more and 10000
seconds or less,
Ms = 519 - 474 × [% C] - 30.4 × [% Mn] - 12.1 × [% Cr] - 7.5 × [% Mo] - 17.7 × [%
Ni] - T1/80
wherein [% C], [% Mn], [% Cr], [% Mo], and [% Ni] indicate the contents (mass%) of
C, Mn, Cr, Mo, and Ni, respectively, and are zero when the element is absent.
- [5] The method for manufacturing the high strength steel sheet according to [4], further
including performing a coating treatment.
Advantageous Effects of Invention
[0011] According to the present invention, a high strength steel sheet can be obtained that
has 980 MPa or higher TS and excels in press formability, flatness in the width direction
and working embrittlement resistance. Furthermore, for example, the high strength
steel sheet of the present invention may be applied to automobile structural members
to reduce the weight of automobile bodies and thereby to enhance fuel efficiency.
Thus, the present invention is highly valuable in industry.
Brief Description of Drawings
[0012]
[Fig. 1] Fig. 1 is a set of views illustrating a structure of a packet having the
largest area in a prior austenite grain according to the present invention, and how
the calculation is made.
[Fig. 2] Fig. 2 is a set of views illustrating the concept of the steepness θ of a
steel sheet according to the present invention, and how the steepness is calculated.
Description of Embodiments
[0013] Embodiments of the present invention will be described below.
[0014] First, appropriate ranges of the chemical composition of the high strength steel
sheet and the reasons why the chemical composition is thus limited will be described.
In the following description, "%" indicating the contents of constituent elements
of steel means "mass%" unless otherwise specified.
[C: 0.030% or more and 0.500% or less]
[0015] Carbon is one of the important basic components of steel. Particularly in the present
invention, carbon is an important element that affects the fraction of martensite
and the working embrittlement resistance. When the C content is less than 0.030%,
the fraction of martensite is so small that realizing 980 MPa or higher TS is difficult.
When, on the other hand, the C content is more than 0.500%, martensite becomes brittle
to cause deterioration in working embrittlement resistance. Thus, the C content is
limited to 0.030% or more and 0.500% or less. The C content is preferably 0.050% or
more. The C content is preferably 0.400% or less. The C content is more preferably
0.100% or more. The C content is more preferably 0.350% or less.
[Si: 0.01% or more and 2.50% or less]
[0016] Silicon is one of the important basic components of steel. Silicon suppresses the
occurrence of carbides during continuous annealing and promotes the formation of retained
austenite. Thus, particularly in the present invention, silicon is an important element
that affects TS and the amount of retained austenite. When the Si content is less
than 0.01%, realizing 980 MPa or higher TS is difficult. When, on the other hand,
the Si content is more than 2.50%, the amount of retained austenite is increased excessively
to make it difficult to achieve 85% or more YR. Thus, the Si content is limited to
0.01% or more and 2.50% or less. The Si content is preferably 0.05% or more. The Si
content is preferably 2.00% or less. The Si content is more preferably 0.10% or more.
The Si content is more preferably 1.20% or less.
[Mn: 0.10% or more and 5.00% or less]
[0017] Manganese is one of the important basic components of steel. Particularly in the
present invention, manganese is an important element that affects the fraction of
martensite and the working embrittlement resistance. When the Mn content is less than
0.10%, the fraction of martensite is so small that realizing 980 MPa or higher TS
is difficult. When, on the other hand, the Mn content is more than 5.00%, martensite
becomes brittle to cause deterioration in working embrittlement resistance. Thus,
the Mn content is limited to 0.10% or more and 5.00% or less. The Mn content is preferably
0.50% or more. The Mn content is preferably 4.50% or less. The Mn content is more
preferably 0.80% or more. The Mn content is more preferably 4.00% or less.
[P: 0.100% or less]
[0018] Phosphorus is segregated at prior austenite grain boundaries and makes the grain
boundaries brittle, thereby lowering the ultimate deformability of steel sheets and
causing deterioration in working embrittlement resistance. Thus, the P content needs
to be 0.100% or less. The lower limit of the P content is not particularly specified.
In view of the fact that phosphorus is a solid solution strengthening element and
can increase the strength of steel sheets, the lower limit is preferably 0.001% or
more. For the reasons above, the P content is limited to 0.100% or less. The P content
is preferably 0.001% or more. The P content is preferably 0.070% or less.
[S: 0.0200% or less]
[0019] Sulfur forms sulfides and lowers the ultimate deformability of steel sheets to cause
deterioration in working embrittlement resistance. Thus, the S content needs to be
0.0200% or less. The lower limit of the S content is not particularly specified but
is preferably 0.0001% or more due to production technique limitations. For the reasons
above, the S content is limited to 0.0200% or less. The S content is preferably 0.0001%
or more. The S content is preferably 0.0050% or less.
[Al: 1.000% or less]
[0020] Aluminum raises the A
3 transformation temperature to allow more ferrite to be contained in the microstructure.
The fraction of martensite is correspondingly lowered to make it difficult to realize
980 MPa or higher TS. Thus, the Al content needs to be 1.000% or less. The lower limit
of the Al content is not particularly specified. In view of the fact that aluminum
suppresses the occurrence of carbides during continuous annealing and promotes the
formation of retained austenite, the Al content is preferably 0.001% or more. For
the reasons above, the Al content is limited to 1.000% or less. The Al content is
preferably 0.001% or more. The Al content is preferably 0.500% or less.
[N: 0.0100% or less]
[0021] Nitrogen forms nitrides and lowers the ultimate deformability of steel sheets to
cause deterioration in working embrittlement resistance. Thus, the N content needs
to be 0.0100% or less. The lower limit of the N content is not particularly specified
but the N content is preferably 0.0001% or more due to production technique limitations.
For the reasons above, the N content is limited to 0.0100% or less. The N content
is preferably 0.0001% or more. The N content is preferably 0.0050% or less.
[O: 0.0100% or less]
[0022] Oxygen forms oxides and lowers the ultimate deformability of steel sheets to cause
deterioration in working embrittlement resistance. Thus, the O content needs to be
0.0100% or less. The lower limit of the O content is not particularly specified but
the O content is preferably 0.0001% or more due to production technique limitations.
For the reasons above, the O content is limited to 0.0100% or less. The O content
is preferably 0.0001% or more. The O content is preferably 0.0050% or less.
[0023] The chemical composition of the high strength steel sheet according to an embodiment
of the present invention includes the components described above, and the balance
is Fe and incidental impurities. Here, the incidental impurities include Zn, Pb, As,
Ge, Sr, and Cs. A total of 0.100% or less of these impurities is acceptable.
[0024] In addition to the components in the proportions described above, the high strength
steel sheet of the present invention may further include at least one element selected
from, in mass%, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10%
or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less,
Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb:
0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr:
0.100% or less, Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less. These
elements may be contained singly or in combination.
[0025] When the contents of Ti, Nb, and V are each 0.200% or less, coarse precipitates and
inclusions will not occur in large amounts and thus will not cause lowering of the
ultimate deformability of steel sheets; hence there will be no deterioration in working
embrittlement resistance. Thus, the contents of Ti, Nb, and V are each preferably
0.200% or less. The lower limits of the contents of Ti, Nb, and V are not particularly
specified. These elements form fine carbides, nitrides, or carbonitrides during hot
rolling or continuous annealing to increase the strength of steel sheets. In view
of this fact, the contents of Ti, Nb, and V are each more preferably 0.001% or more.
When titanium, niobium, and vanadium are added, the contents thereof are each limited
to 0.200% or less for the reasons above. The contents are each more preferably 0.001%
or more. The contents are each more preferably 0.100% or less.
[0026] When the contents of Ta and W are each 0.10% or less, coarse precipitates and inclusions
will not occur in large amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the contents of Ta and W are each preferably 0.10% or less. The lower limits
of the contents of Ta and W are not particularly specified. These elements form fine
carbides, nitrides, or carbonitrides during hot rolling or continuous annealing to
increase the strength of steel sheets in some cases. In view of this fact, the contents
of Ta and W are each more preferably 0.01% or more. When tantalum and tungsten are
added, the contents thereof are each limited to 0.10% or less for the reasons above.
The contents are each more preferably 0.01% or more. The contents are each more preferably
0.08% or less.
[0027] When the B content is 0.0100% or less, inner cracks that lower the ultimate deformability
of steel sheets will not form during casting or hot rolling and thus there will be
no deterioration in working embrittlement resistance. Thus, the B content is preferably
0.0100% or less. The lower limit of the B content is not particularly specified. The
B content is more preferably 0.0003% or more in view of the fact that this element
is segregated at austenite grain boundaries during annealing and enhances hardenability.
When boron is added, the content thereof is limited to 0.0100% or less for the reasons
above. The content is more preferably 0.0003% or more. The content is more preferably
0.0080% or less.
[0028] When the contents of Cr, Mo, and Ni are each 1.00% or less, coarse precipitates and
inclusions will not occur in increased amounts and thus will not cause lowering of
the ultimate deformability of steel sheets; hence there will be no deterioration in
working embrittlement resistance. Thus, the contents of Cr, Mo, and Ni are each preferably
1.00% or less. The lower limits of the contents of Cr, Mo, and Ni are not particularly
specified. In view of the fact that these elements enhance hardenability, the contents
of Cr, Mo, and Ni are each more preferably 0.01% or more. When chromium, molybdenum,
and nickel are added, the contents thereof are each limited to 1.00% or less for the
reasons above. The contents are each more preferably 0.01% or more. The contents are
each more preferably 0.80% or less.
[0029] When the Co content is 0.010% or less, coarse precipitates and inclusions will not
occur in increased amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the Co content is preferably 0.010% or less. The lower limit of the Co content
is not particularly specified. In view of the fact that this element enhances hardenability,
the Co content is more preferably 0.001% or more. When cobalt is added, the content
thereof is limited to 0.010% or less for the reasons above. The content is more preferably
0.001% or more. The content is more preferably 0.008% or less.
[0030] When the Cu content is 1.00% or less, coarse precipitates and inclusions will not
occur in increased amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the Cu content is preferably 1.00% or less. The lower limit of the Cu content
is not particularly specified. In view of the fact that this element enhances hardenability,
the Cu content is preferably 0.01% or more. When copper is added, the content thereof
is limited to 1.00% or less for the reasons above. The content is more preferably
0.01% or more. The content is more preferably 0.80% or less.
[0031] When the Sn content is 0.200% or less, inner cracks that lower the ultimate deformability
of steel sheets will not form during casting or hot rolling and thus there will be
no deterioration in working embrittlement resistance. Thus, the Sn content is preferably
0.200% or less. The lower limit of the Sn content is not particularly specified. The
Sn content is more preferably 0.001% or more in view of the fact that tin enhances
hardenability (in general, is an element that enhances corrosion resistance). When
tin is added, the content thereof is limited to 0.200% or less for the reasons above.
The content is more preferably 0.001% or more. The content is more preferably 0.100%
or less.
[0032] When the Sb content is 0.200% or less, coarse precipitates and inclusions will not
occur in increased amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the Sb content is preferably 0.200% or less. The lower limit of the Sb content
is not particularly specified. In view of the fact that this element enables control
of the thickness of surface layer softening and the strength, the Sb content is more
preferably 0.001% or more. When antimony is added, the content thereof is limited
to 0.200% or less for the reasons above. The content is more preferably 0.001% or
more. The content is more preferably 0.100% or less.
[0033] When the contents of Ca, Mg, and REM are each 0.0100% or less, coarse precipitates
and inclusions will not occur in increased amounts and thus will not cause lowering
of the ultimate deformability of steel sheets; hence there will be no deterioration
in working embrittlement resistance. Thus, the contents of Ca, Mg, and REM are each
preferably 0.0100% or less. The lower limits of the contents of Ca, Mg, and REM are
not particularly specified. In view of the fact that these elements change the shapes
of nitrides and sulfides into spheroidal and enhance the ultimate deformability of
steel sheets, the contents of Ca, Mg, and REM are each more preferably 0.0005% or
more. When calcium, magnesium, and rare earth metal(s) are added, the contents thereof
are each limited to 0.0100% or less for the reasons above. The contents are each more
preferably 0.0005% or more. The contents are each more preferably 0.0050% or less.
[0034] When the contents of Zr and Te are each 0.100% or less, coarse precipitates and inclusions
will not occur in increased amounts and thus will not cause lowering of the ultimate
deformability of steel sheets; hence there will be no deterioration in working embrittlement
resistance. Thus, the contents of Zr and Te are each preferably 0.100% or less. The
lower limits of the contents of Zr and Te are not particularly specified. In view
of the fact that these elements change the shapes of nitrides and sulfides into spheroidal
and enhance the ultimate deformability of steel sheets, the contents of Zr and Te
are each more preferably 0.001% or more. When zirconium and tellurium are added, the
contents thereof are each limited to 0.100% or less for the reasons above. The contents
are each more preferably 0.001% or more. The contents are each more preferably 0.080%
or less.
[0035] When the Hf content is 0.10% or less, coarse precipitates and inclusions will not
occur in increased amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the Hf content is preferably 0.10% or less. The lower limit of the Hf content
is not particularly specified. In view of the fact that this element changes the shapes
of nitrides and sulfides into spheroidal and enhances the ultimate deformability of
steel sheets, the Hf content is more preferably 0.01% or more. When hafnium is added,
the content thereof is limited to 0.10% or less for the reasons above. The content
is more preferably 0.01% or more. The content is more preferably 0.08% or less.
[0036] When the Bi content is 0.200% or less, coarse precipitates and inclusions will not
occur in increased amounts and thus will not cause lowering of the ultimate deformability
of steel sheets; hence there will be no deterioration in working embrittlement resistance.
Thus, the Bi content is preferably 0.200% or less. The lower limit of the Bi content
is not particularly specified. In view of the fact that this element reduces the occurrence
of segregation, the Bi content is more preferably 0.001% or more. When bismuth is
added, the content thereof is limited to 0.200% or less for the reasons above. The
content is more preferably 0.001% or more. The content is more preferably 0.100% or
less.
[0037] When the content of any of Ti, Nb, V, Ta, W, B, Cr, Mo, Ni, Co, Cu, Sn, Sb, Ca, Mg,
REM, Zr, Te, Hf, and Bi is below the preferred lower limit, the element does not impair
the advantageous effects of the present invention and is regarded as an incidental
impurity.
[0038] Next, the steel microstructure of the high strength steel sheet of the present invention
will be described.
[Area fraction of tempered martensite: 38% or more and less than 90%]
[0039] When the amount of tempered martensite is less than 38%, realizing 980 MPa or higher
TS is difficult. When, on the other hand, the amount of tempered martensite is 90%
or more, the amount of ferrite is lowered to cause a decrease in El and consequently
press formability is lowered. Thus, the amount of tempered martensite is limited to
38% or more and less than 90%. The amount is preferably 40% or more. The amount is
preferably 60% or less.
[0040] Here, tempered martensite is measured as follows. A longitudinal cross section of
the steel sheet is polished and is etched with 3 vol% Nital. A portion at 1/4 sheet
thickness (a location corresponding to 1/4 of the sheet thickness in the depth direction
from the steel sheet surface) is observed using SEM in 10 fields of view at a magnification
of ×2000. In the microstructure images, tempered martensite is structures that have
fine irregularities inside the structures and contain inner carbides. The values thus
obtained are averaged to determine the tempered martensite.
[Amount of retained austenite: less than 3%]
[0041] This configuration is a very important requirement that constitutes the present invention.
When the volume fraction of retained austenite is 3% or more, press formability is
lowered. The reason for low press formability is that retained austenite with a high
fraction gives rise to a lowering in λ by undergoing strain-induced transformation.
Thus, the retained austenite is limited to less than 3%. The amount of retained austenite
is preferably 1% or less. The lower limit of retained austenite is not particularly
limited and may be 0%.
[0042] Here, retained austenite is measured as follows. The steel sheet is polished to expose
a face 0.1 mm below 1/4 sheet thickness and is thereafter further chemically polished
to expose a face 0.1 mm below the face exposed above. The face is analyzed with an
X-ray diffractometer using CoKα radiation to determine the integral intensity ratios
of the diffraction peaks of {200}, {220}, and {311} planes of fcc iron and {200},
{211}, and {220} planes of bcc iron. Nine integral intensity ratios thus obtained
are averaged to determine retained austenite.
[Area fraction of the total of ferrite and bainitic ferrite: 10% or more and 60% or
less]
[0043] This configuration is a very important requirement that constitutes the present invention.
When the total of ferrite and bainitic ferrite is less than 10%, El is lowered and
consequently press formability is deteriorated. When, on the other hand, the total
of ferrite and bainitic ferrite is more than 60%, realizing 980 MPa or higher TS is
difficult. Thus, the total of ferrite and bainitic ferrite is limited to 10% or more
and 60% or less. The total amount is preferably 35% or more. The total amount is preferably
55% or less.
[0044] Here, the total of ferrite and bainitic ferrite is measured as follows. A longitudinal
cross section of the steel sheet is polished and is etched with 3 vol% Nital. A portion
at 1/4 sheet thickness (a location corresponding to 1/4 of the sheet thickness in
the depth direction from the steel sheet surface) is observed using SEM in 10 fields
of view at a magnification of ×2000. In the microstructure images, ferrite is recessed
structures having a flat interior and containing no inner carbides. In the microstructure
images, bainitic ferrite is recessed structures having a flat interior and containing
inner carbides. The values thus obtained are combined and are averaged to determine
the total of ferrite and bainitic ferrite.
[0045] Possible microstructures other than those described above include pearlite, fresh
martensite, and acicular ferrite. These microstructures do not affect characteristics
as long as their fractions are 5% or less, and thus may be present within that range.
[Average grain size of prior austenite: 20 um or less]
[0046] This configuration is a very important requirement that constitutes the present invention.
Reducing the average grain size of prior austenite can suppress crack propagation
and thereby enhances the working embrittlement resistance of steel sheets. In order
to obtain these effects, the average grain size of prior austenite needs to be 20
um or less. The lower limit of the average grain size of prior austenite is not particularly
specified. When, however, the average grain size of prior austenite is less than 2
um, more retained austenite may form. Thus, the average grain size is preferably 2
um or more. For the reasons above, the average grain size of prior austenite is limited
to 20 um or less. The average grain size is preferably 2 um or more. The average grain
size is preferably 15 um or less. The average grain size is more preferably 3 um or
more. The average grain size is more preferably 10 um or less.
[0047] Here, the average grain size of prior austenite is measured as follows. A longitudinal
cross section of the steel sheet is polished and is etched with, for example, a mixed
solution of picric acid and ferric chloride to expose prior austenite grain boundaries.
Portions at 1/4 sheet thickness (locations corresponding to 1/4 of the sheet thickness
in the depth direction from the steel sheet surface) are photographed with an optical
microscope each in 3 to 10 fields of view at a magnification of ×400. Twenty straight
lines including 10 vertical lines and 10 horizontal lines are drawn at regular intervals
on the image data obtained, and the grain size is determined by a linear intercept
method.
[Average of the proportions of packets having the largest area in prior austenite
grains: 70% by area or less]
[0048] This configuration is a very important requirement that constitutes the present invention.
The proportion of a packet having the largest area in a prior austenite grain affects
the flatness in the width direction and the working embrittlement resistance. As illustrated
in Fig. 1, a prior austenite grain contains up to four kinds of packets distinguished
by crystal habit plane formed by transformation. The packet having the largest area
in a prior austenite grain is the packet that occupies the largest area among such
packets.
[0049] The proportion of one packet in a prior austenite grain is determined by dividing
the area of the packet of interest by the area of the whole prior austenite grain.
As a result of extensive studies, the present inventors have found that strain among
the packets is reduced and the flatness in the width direction is improved by lowering
the proportion of a packet having the largest area in a prior austenite grain. The
present inventors have also found that lowering the proportion of a packet having
the largest area in a prior austenite grain leads to a fine microstructure and suppresses
crack propagation, thereby enhancing the working embrittlement resistance of the steel
sheet. Thus, the average of the proportions of packets having the largest area in
prior austenite grains is limited to 70% or less. The average proportion is preferably
60% or less. The lower limit of the average proportion of packets having the largest
area in prior austenite grains is not particularly limited. The grains contain up
to four kinds of packets. When four packets are evenly distributed, the proportion
of a packet having the largest area in the prior austenite grain is 25%. Thus, the
lower limit of the average proportion of packets having the largest area in prior
austenite grains may be 25% or more but is not necessarily limited thereto.
[0050] Here, the average proportion of packets having the largest area in prior austenite
grains is measured as follows. First, a test specimen for microstructure observation
is sampled from the cold rolled steel sheet. Next, the sampled test specimen is polished
by vibration polishing with colloidal silica to expose a cross section in the rolling
direction (a longitudinal cross section) for use as observation surface. The observation
surface is specular. Next, electron backscatter diffraction (EBSD) measurement is
performed with respect to a portion at 1/4 sheet thickness (a location corresponding
to 1/4 of the sheet thickness in the depth direction from the steel sheet surface)
to obtain local crystal orientation data. Here, the SEM magnification is ×1000, the
step size is 0.2 um, the measured region is 80 um square, and the WD is 15 mm. The
local orientation data obtained is analyzed with OIM Analysis 7 (OIM), and a map (a
CP map) that shows close-packed plane groups (CP groups) with different colors is
created using the method described in Non Patent Literature 1. In the present invention,
a packet is defined as a region or regions belonging to the same CP group. From the
CP map obtained, the area of the packet having the largest area is determined and
is divided by the area of the whole prior austenite grain to give the proportion of
the packet having the largest area in the prior austenite grain. This analysis is
performed with respect to 10 or more adjacent prior austenite grains, and the results
are averaged to give the average proportion of packets having the largest area in
prior austenite grains.
[0051] Next, a manufacturing method of the present invention will be described.
[0052] In the present invention, a steel material (a steel slab) may be obtained by any
known steelmaking method without limitation, such as a converter or an electric arc
furnace. To prevent macro-segregation, the steel slab (the slab) is preferably produced
by a continuous casting method.
[0053] In the present invention, the slab heating temperature, the slab soaking holding
time, and the coiling temperature in hot rolling are not particularly limited. For
example, the steel slab may be hot rolled in such a manner that the slab is heated
and is then rolled, that the slab is subjected to hot direct rolling after continuous
casting without being heated, or that the slab is subjected to a short heat treatment
after continuous casting and is then rolled. The slab heating temperature, the slab
soaking holding time, the finish rolling temperature, and the coiling temperature
in hot rolling are not particularly limited. The lower limit of the slab heating temperature
is preferably 1100°C or above. The upper limit of the slab heating temperature is
preferably 1300°C or below. The lower limit of the slab soaking holding time is preferably
30 minutes or more. The upper limit of the slab soaking holding time is preferably
250 minutes or less. The lower limit of the finish rolling temperature is preferably
Ar
3 transformation temperature or above. Furthermore, the lower limit of the coiling
temperature is preferably 350°C or above. The upper limit of the coiling temperature
is preferably 650°C or below.
[0054] The hot rolled steel sheet thus produced is pickled. Pickling can remove oxides on
the steel sheet surface and is thus important to ensure good chemical convertibility
and a high quality of coating in the final high strength steel sheet. Pickling may
be performed at a time or several. The hot rolled sheet that has been pickled may
be cold rolled directly or may be subjected to heat treatment before cold rolling.
[0055] The rolling reduction in cold rolling and the sheet thickness after rolling are not
particularly limited. The lower limit of the rolling reduction is preferably 30% or
more. The upper limit of the rolling reduction is preferably 80% or less. The advantageous
effects of the present invention may be obtained without any limitations on the number
of rolling passes and the rolling reduction in each pass.
[0056] The cold rolled steel sheet obtained as described above is annealed. Annealing conditions
are as follows.
[Annealing temperature T1: 700°C or above and 950°C or below]
[0057] When the annealing temperature T1 is below 700°C, the area fraction of the total
of ferrite and bainitic ferrite is more than 60% to make it difficult to realize 980
MPa or higher TS. When, on the other hand, the annealing temperature T1 is above 950°C,
prior austenite grains are excessively increased in size and the prior austenite grain
size exceeds 20 um to give rise to a decrease in working embrittlement resistance.
Thus, the annealing temperature T1 is limited to 700°C or above and 950°C or below.
The annealing temperature T1 is preferably 800°C or above. The annealing temperature
T1 is preferably 900°C or below.
[Holding time t1 at the annealing temperature T1: 10 seconds or more and 1000 seconds
or less]
[0058] When the holding time t1 at the annealing temperature T1 is less than 10 seconds,
austenitization is insufficient and the area fraction of the total of ferrite and
bainitic ferrite is more than 60%. As a result, it is difficult to achieve 980 MPa
or higher TS. When, on the other hand, the holding time at the annealing temperature
T1 is more than 1000 seconds, the prior austenite grain size is excessively increased,
and the working embrittlement resistance is lowered. For the reasons above, the holding
time t1 at the annealing temperature T1 is limited to 10 seconds or more and 1000
seconds or less. The holding time t1 is preferably 50 seconds or more. The holding
time t1 is preferably 500 seconds or less.
[Average cooling rate from 750°C to 600°C: less than 20°C/s]
[0059] When the average cooling rate from 750°C to 600°C is 20°C/s or more, the area fraction
of the total of ferrite and bainitic ferrite is less than 10% to cause a decrease
in El, thereby deteriorating press formability. For the reasons above, the average
cooling rate from 750°C to 600°C is limited to less than 20°C/s. The average cooling
rate is preferably 15°C/s or less.
[Average cooling rate from (Ms + 50°C) to a quench start temperature T2: less than
5°C/s]
[0060] This configuration is a very important requirement that constitutes the present invention.
When the average cooling rate from (Ms + 50°C) to a quench start temperature T2 is
5°C/s or more, the area fraction of the total of ferrite and bainitic ferrite is less
than 10% to cause a decrease in El, thereby deteriorating press formability. For the
reasons above, the average cooling rate from (Ms + 50°C) to a quench start temperature
T2 is limited to less than 5°C/s. The average cooling rate is preferably 4°C/s or
less.
[Quench start temperature T2: (Ms - 80°C) or above and below Ms]
[0061] This configuration is a very important requirement that constitutes the present invention.
The quench start temperature T2 is controlled to (Ms - 80°C) or above and below Ms
to ensure that the martensite transformation rate before the start of quenching is
1% or more and 80% or less. In this manner, quenching can give microstructures in
which the average proportion of packets having the largest area in prior austenite
grains is 70% or less and the volume fraction of retained austenite is less than 3%.
When the quench start temperature T2 is below (Ms - 80°C), the martensite transformation
rate before the start of quenching exceeds 80% and consequently the volume fraction
of retained austenite is 3% or more to cause a decrease in press formability. When,
on the other hand, the quench start temperature T2 is above Ms, the martensite transformation
rate before the start of quenching is less than 1% and the average proportion of packets
having the largest area in prior austenite grains exceeds 70% to cause deterioration
in flatness in the width direction and working embrittlement resistance. Thus, the
quench start temperature T2 is limited to (Ms - 80°C) or above and below Ms. The quench
start temperature T2 is preferably (Ms - 50°C) or above. The quench start temperature
T2 is preferably (Ms - 5°C) or below. The martensite start temperature Ms (°C) is
defined by the following formula (1):
Ms = 519 - 474 × [% C] - 30.4 × [% Mn] - 12.1 × [% Cr] - 7.5 × [% Mo] - 17.7 × [%
Ni] - T1/80
wherein [% C], [% Mn], [% Cr], [% Mo], and [% Ni] indicate the contents (mass%) of
C, Mn, Cr, Mo, and Ni, respectively, and are zero when the element is absent.
[Average cooling rate from the quench start temperature T2 to 80°C: 300°C/s or more]
[0062] When the average cooling rate from the quench start temperature T2 to 80°C is less
than 300°C/s, the volume fraction of retained austenite is 3% or more to cause a decrease
in press formability. Thus, the average cooling rate from the quench start temperature
T2 to 80°C is limited to 300°C/s or more. The average cooling rate is preferably 800°C/s
or more. The upper limit is not necessarily specified but is preferably 2000°C/s or
less.
[Tempering temperature T3: 100°C or above and 400°C or below]
[0063] In the present invention, tempered martensite is a microstructure that is formed
when martensite at 80°C or below is heat-treated at a tempering temperature of 100°C
or above for a holding time of 10 seconds or more. Thus, martensite is not sufficiently
tempered when the tempering temperature T3 is below 100°C. The resultant microstructures
will be based on as-quenched martensite, which deteriorates the working embrittlement
resistance. When, on the other hand, the tempering temperature T3 is above 400°C,
martensite is excessively tempered to make it difficult to achieve 980 MPa or higher
TS. For the reasons above, the tempering temperature T3 is limited to 100°C or above
and 400°C or below. The tempering temperature T3 is preferably 150°C or above. The
tempering temperature T3 is preferably 350°C or below.
[Holding time t3 at the tempering temperature T3: 10 seconds or more and 10000 seconds
or less]
[0064] In the present invention, tempered martensite is a microstructure that is formed
when martensite at 80°C or below is heat-treated at a tempering temperature of 100°C
or above for a holding time of 10 seconds or more. Thus, martensite is not sufficiently
tempered when the holding time t3 at the tempering temperature T3 is less than 10
seconds. The resultant microstructures will be based on as-quenched martensite, which
deteriorates the working embrittlement resistance. When, on the other hand, the tempering
temperature T3 is more than 10000 seconds, martensite is excessively tempered to make
it difficult to achieve 980 MPa or higher TS. For the reasons above, the holding time
t3 at the tempering temperature T3 is limited to 10 seconds or more and 10000 seconds
or less. The holding time t3 is preferably 50 seconds or more. The holding time t3
is preferably 5000 seconds or less.
[0065] Post-temper cooling is not particularly limited and the steel sheet may be cooled
to a desired temperature in an appropriate manner. Incidentally, the desired temperature
is preferably about room temperature.
[0066] Furthermore, the high strength steel sheet described above may be worked under conditions
where the amount of equivalent plastic strain is 0.10% or more and 5.00% or less.
The working may be followed by reheating at 100°C or above and 400°C or below.
[0067] When the high strength steel sheet is a product that is traded, the steel sheet is
usually traded after being cooled to room temperature.
[0068] The high strength steel sheet may be subjected to coating treatment during annealing
or after annealing.
[0069] For example, the coating treatment during annealing may be hot-dip galvanizing treatment
performed when the steel sheet is being cooled or has been cooled from 750°C to 600°C
at an average cooling rate of less than 20°C/s. The hot-dip galvanizing treatment
may be followed by alloying. For example, the coating treatment after annealing may
be Zn-Ni electrical alloy coating treatment or pure Zn electroplated coating treatment
performed after tempering. A coated layer may be formed by electroplated coating,
or hot-dip zinc-aluminum-magnesium alloy coating may be applied. While the coating
treatment has been described above focusing on zinc coating, the types of coating
metals, such as Zn coating and Al coating, are not particularly limited. Other conditions
in the manufacturing method are not particularly limited. From the point of view of
productivity, the series of treatments including annealing, hot-dip galvanizing, and
alloying treatment of the coated zinc layer is preferably performed on hot-dip galvanizing
line CGL (continuous galvanizing line). To control the coating weight of the coated
layer, the hot-dip galvanizing treatment may be followed by wiping. Conditions for
operations, such as coating, other than those conditions described above may be determined
in accordance with the usual hot-dip galvanizing technique.
[0070] After the coating treatment after annealing, the steel sheet may be worked again
under conditions where the amount of equivalent plastic strain is 0.10% or more and
5.00 or less. The working may be followed by reheating at 100°C or above and 400°C
or below.
EXAMPLES
[0071] Steels having a chemical composition described in Table 1 and 2, with the balance
being Fe and incidental impurities, were smelted in a converter and were continuously
cast into slabs. Next, the slabs obtained were heated, hot rolled, pickled, cold rolled,
and subjected to annealing treatment and tempering treatment described in Tables 3
to 5. High strength cold rolled steel sheets having a sheet thickness of 0.6 to 2.2
mm were thus obtained. Incidentally, some of the steel sheets were subjected to coating
treatment during or after annealing.
[Table 1]
Steels |
Chemical composition (mass%) |
|
C |
Si |
Mn |
P |
S |
N |
O |
Al |
Ti |
B |
Nb |
Cu |
Others |
A |
0.215 |
0.280 |
2.24 |
0.006 |
0.0011 |
0.005 |
0.006 |
0.047 |
|
|
|
|
|
INV. EX. |
B |
0.217 |
0.298 |
1.98 |
0.005 |
0.0009 |
0.006 |
0.005 |
0.053 |
|
|
|
|
|
INV. EX. |
C |
0.192 |
0.251 |
2.04 |
0.009 |
0.0008 |
0.002 |
0.003 |
0.015 |
|
|
|
|
|
INV. EX. |
D |
0.111 |
1.332 |
2.04 |
0.011 |
0.0007 |
0.004 |
0.005 |
0.024 |
|
|
|
|
|
INV. EX. |
E |
0.113 |
1.464 |
2.22 |
0.010 |
0.0014 |
0.006 |
0.002 |
0.018 |
|
|
|
|
|
INV. EX. |
F |
0.048 |
0.262 |
2.25 |
0.008 |
0.0011 |
0.004 |
0.002 |
0.030 |
|
|
|
|
|
INV. EX. |
G |
0.021 |
0.168 |
2.30 |
0.007 |
0.0013 |
0.003 |
0.005 |
0.036 |
|
|
|
|
|
COMP. EX. |
H |
0.468 |
0.166 |
2.09 |
0.009 |
0.0012 |
0.006 |
0.002 |
0.059 |
|
|
|
|
|
INV. EX. |
I |
0.522 |
0.248 |
2.00 |
0.010 |
0.0015 |
0.001 |
0.007 |
0.015 |
|
|
|
|
|
COMP. EX. |
J |
0.212 |
0.073 |
2.13 |
0.015 |
0.0010 |
0.005 |
0.006 |
0.052 |
|
|
|
|
|
INV. EX. |
K |
0.191 |
0.002 |
2.23 |
0.013 |
0.0006 |
0.002 |
0.006 |
0.020 |
|
|
|
|
|
COMP. EX. |
L |
0.210 |
2.339 |
2.22 |
0.006 |
0.0010 |
0.006 |
0.007 |
0.054 |
|
|
|
|
|
INV. EX. |
M |
0.205 |
2.532 |
1.98 |
0.012 |
0.0007 |
0.002 |
0.006 |
0.025 |
|
|
|
|
|
COMP. EX. |
N |
0.203 |
0.273 |
0.27 |
0.006 |
0.0007 |
0.001 |
0.004 |
0.055 |
|
|
|
|
|
INV. EX. |
O |
0.187 |
0.307 |
0.08 |
0.012 |
0.0009 |
0.007 |
0.005 |
0.033 |
|
|
|
|
|
COMP. EX. |
P |
0.191 |
0.271 |
4.98 |
0.012 |
0.0011 |
0.005 |
0.002 |
0.052 |
|
|
|
|
|
INV. EX. |
Q |
0.197 |
0.162 |
5.12 |
0.007 |
0.0007 |
0.005 |
0.004 |
0.040 |
|
|
|
|
|
COMP. EX. |
R |
0.216 |
0.314 |
2.12 |
0.099 |
0.0005 |
0.002 |
0.003 |
0.012 |
|
|
|
|
|
INV. EX. |
S |
0.219 |
0.333 |
2.06 |
0.121 |
0.0010 |
0.004 |
0.003 |
0.049 |
|
|
|
|
|
COMP. EX. |
T |
0.219 |
0.173 |
2.25 |
0.006 |
0.0182 |
0.006 |
0.006 |
0.024 |
|
|
|
|
|
INV. EX. |
U |
0.205 |
0.192 |
2.06 |
0.010 |
0.0222 |
0.004 |
0.002 |
0.024 |
|
|
|
|
|
COMP. EX. |
V |
0.205 |
0.165 |
2.18 |
0.011 |
0.0010 |
0.002 |
0.003 |
0.976 |
|
|
|
|
|
INV. EX. |
W |
0.195 |
0.165 |
2.10 |
0.015 |
0.0009 |
0.005 |
0.002 |
1.135 |
|
|
|
|
|
COMP. EX. |
X |
0.214 |
0.204 |
2.06 |
0.011 |
0.0009 |
0.0089 |
0.007 |
0.059 |
|
|
|
|
|
INV. EX. |
Y |
0.192 |
0.186 |
1.93 |
0.012 |
0.0007 |
0.0112 |
0.002 |
0.032 |
|
|
|
|
|
COMP. EX. |
Z |
0.182 |
0.229 |
2.22 |
0.011 |
0.0007 |
0.004 |
0.0090 |
0.037 |
|
|
|
|
|
INV. EX. |
AA |
0.206 |
0.271 |
2.19 |
0.009 |
0.0015 |
0.006 |
0.0110 |
0.040 |
|
|
|
|
|
COMP. EX. |
AB |
0.186 |
0.331 |
1.91 |
0.007 |
0.0010 |
0.005 |
0.001 |
0.038 |
0.002 |
|
|
|
|
INV. EX. |
AC |
0.190 |
0.327 |
2.29 |
0.013 |
0.0008 |
0.006 |
0.002 |
0.027 |
0.187 |
|
|
|
|
INV. EX. |
AD |
0.206 |
0.341 |
2.26 |
0.011 |
0.0012 |
0.006 |
0.003 |
0.038 |
0.223 |
|
|
|
|
COMP. EX. |
AE |
0.182 |
0.258 |
2.06 |
0.008 |
0.0009 |
0.005 |
0.005 |
0.038 |
|
0.0002 |
|
|
|
INV. EX. |
AF |
0.189 |
0.309 |
2.12 |
0.006 |
0.0005 |
0.007 |
0.003 |
0.048 |
|
0.0088 |
|
|
|
INV. EX. |
AG |
0.206 |
0.239 |
2.16 |
0.008 |
0.0009 |
0.003 |
0.005 |
0.031 |
|
0.0121 |
|
|
|
COMP. EX. |
AH |
0.187 |
0.164 |
2.27 |
0.006 |
0.0013 |
0.002 |
0.004 |
0.016 |
|
|
0.002 |
|
|
INV. EX. |
AI |
0.216 |
0.308 |
2.20 |
0.008 |
0.0011 |
0.002 |
0.003 |
0.031 |
|
|
0.189 |
|
|
INV. EX. |
AJ |
0.190 |
0.189 |
2.05 |
0.012 |
0.0006 |
0.003 |
0.007 |
0.051 |
|
|
0.203 |
|
|
COMP. EX. |
AK |
0.183 |
0.345 |
2.03 |
0.009 |
0.0013 |
0.002 |
0.005 |
0.023 |
|
|
|
0.03 |
|
INV. EX. |
Underlines indicate being outside the range of the present invention. |
[Table 2]
Steels |
Chemical composition (mass%) |
|
C |
Si |
Mn |
P |
S |
N |
O |
Al |
Ti |
B |
Nb |
Cu |
Others |
AL |
0.205 |
0.293 |
1.98 |
0.013 |
0.0012 |
0.005 |
0.002 |
0.057 |
|
|
|
0.90 |
|
INV. EX. |
AM |
0.190 |
0.317 |
2.29 |
0.007 |
0.0010 |
0.003 |
0.007 |
0.052 |
|
|
|
1.11 |
|
COMP. EX. |
AN |
0.214 |
0.261 |
2.25 |
0.007 |
0.0010 |
0.005 |
0.004 |
0.050 |
|
|
|
|
V:0.070 |
INV. EX. |
AO |
0.205 |
0.309 |
2.02 |
0.007 |
0.0006 |
0.002 |
0.006 |
0.028 |
|
|
|
|
Ta:0.05 |
INV. EX. |
AP |
0.213 |
0.215 |
2.26 |
0.014 |
0.0006 |
0.006 |
0.002 |
0.048 |
|
|
|
|
W:0.03 |
INV. EX. |
AQ |
0.187 |
0.263 |
2.10 |
0.015 |
0.0007 |
0.005 |
0.003 |
0.036 |
|
|
|
|
Cr:0.87 |
INV. EX. |
AR |
0.184 |
0.228 |
2.06 |
0.009 |
0.0011 |
0.002 |
0.004 |
0.029 |
|
|
|
|
Mo:0.13 |
INV. EX. |
AS |
0.193 |
0.288 |
2.29 |
0.012 |
0.0005 |
0.006 |
0.004 |
0.028 |
|
|
|
|
Co:0.008 |
INV. EX. |
AT |
0.219 |
0.242 |
2.21 |
0.007 |
0.0012 |
0.002 |
0.003 |
0.034 |
|
|
|
|
Ni:0.33 |
INV. EX. |
AU |
0.185 |
0.245 |
1.92 |
0.012 |
0.0011 |
0.007 |
0.004 |
0.056 |
|
|
|
|
Sn:0.012 |
INV. EX. |
AV |
0.185 |
0.301 |
2.08 |
0.009 |
0.0006 |
0.005 |
0.004 |
0.026 |
|
|
|
|
Sb:0.005 |
INV. EX. |
AW |
0.196 |
0.274 |
1.90 |
0.010 |
0.0012 |
0.004 |
0.002 |
0.056 |
|
|
|
|
Ca:0.0015 |
INV. EX. |
AX |
0.218 |
0.339 |
1.95 |
0.009 |
0.0009 |
0.004 |
0.004 |
0.022 |
|
|
|
|
Mg:0.0086 |
INV. EX. |
AY |
0.186 |
0.343 |
2.010 |
0.005 |
0.0008 |
0.002 |
0.002 |
0.051 |
|
|
|
|
Zr:0.083 |
INV. EX. |
AZ |
0.217 |
0.240 |
1.990 |
0.008 |
0.0013 |
0.007 |
0.003 |
0.024 |
|
|
|
|
Te:0.092 |
INV. EX. |
BA |
0.102 |
1.364 |
2.270 |
0.010 |
0.0014 |
0.006 |
0.005 |
0.016 |
|
|
|
|
Hf:0.05 |
INV. EX. |
BB |
0.128 |
1.390 |
1.960 |
0.012 |
0.0005 |
0.005 |
0.004 |
0.037 |
|
|
|
|
REM:0.0092 |
INV. EX. |
BC |
0.137 |
1.402 |
2.020 |
0.012 |
0.0007 |
0.003 |
0.005 |
0.030 |
|
|
|
|
Bi:0.164 |
INV. EX. |
BD |
0.132 |
1.328 |
1.940 |
0.007 |
0.0006 |
0.004 |
0.003 |
0.012 |
|
|
|
|
Zn:0.03 |
INV. EX. |
BE |
0.113 |
1.482 |
2.090 |
0.008 |
0.0007 |
0.005 |
0.007 |
0.041 |
|
|
|
|
Pb:0.016 |
INV. EX. |
BF |
0.111 |
1.374 |
2.000 |
0.007 |
0.0008 |
0.005 |
0.002 |
0.011 |
|
|
|
|
As:0.040 |
INV. EX. |
BG |
0.116 |
1.362 |
2.080 |
0.011 |
0.0011 |
0.006 |
0.006 |
0.012 |
|
|
|
|
Ge:0.090 |
INV. EX. |
BH |
0.133 |
1.387 |
2.220 |
0.009 |
0.0009 |
0.001 |
0.003 |
0.052 |
|
|
|
|
Sr:0.065 |
INV. EX. |
BI |
0.108 |
1.310 |
2.150 |
0.007 |
0.0012 |
0.001 |
0.004 |
0.037 |
|
|
|
|
Cs:0.082 |
INV. EX. |
BJ |
0.198 |
0.870 |
2.700 |
0.010 |
0.0003 |
0.004 |
0.001 |
0.045 |
0.007 |
0.0017 |
0.014 |
0.18 |
Ni:0.05 |
INV. EX. |
BK |
0.218 |
0.326 |
2.060 |
0.009 |
0.0008 |
0.007 |
0.007 |
0.012 |
|
|
|
|
|
INV. EX. |
BL |
0.108 |
1.352 |
1.920 |
0.013 |
0.0011 |
0.003 |
0.004 |
0.056 |
|
|
|
|
|
INV. EX. |
BM |
0.105 |
1.331 |
2.030 |
0.007 |
0.0006 |
0.004 |
0.002 |
0.049 |
|
|
|
|
|
INV. EX. |
BN |
0.207 |
1.374 |
1.910 |
0.007 |
0.0013 |
0.003 |
0.001 |
0.057 |
|
|
|
|
|
INV. EX. |
BO |
0.189 |
1.414 |
2.020 |
0.009 |
0.0015 |
0.003 |
0.003 |
0.050 |
|
|
|
|
|
INV. EX. |
Underlines indicate being outside the range of the present invention. |
[Table 3]
Nos. |
Steels |
Annealing temp. T1 (°C) |
Holding time t1 (s) |
Average cooling rate in temperature range of 750-600°C (°C/s) |
Average cooling rate in temperature range of (Ms+50°C)-quench start temp. T2 (°C/s) |
Ms (°C) |
(Ms-80) (°C) |
Quench start temp. T2 (°C) |
Cooling rate from T2 to 80°C (°C/s) |
Tempering temp. 13 (°C) |
Holding time t3 (s) |
Type" |
|
1 |
A |
796 |
322 |
13 |
2 |
330 |
250 |
321 |
905 |
203 |
854 |
CR |
INV. EX. |
2 |
B |
783 |
348 |
7 |
3 |
337 |
257 |
323 |
966 |
189 |
994 |
CR |
INV. EX. |
3 |
B |
717 |
427 |
12 |
2 |
337 |
257 |
320 |
875 |
172 |
961 |
CR |
INV. EX. |
4 |
B |
692 |
255 |
10 |
2 |
337 |
257 |
327 |
957 |
215 |
989 |
CR |
COMP. EX. |
5 |
B |
927 |
274 |
7 |
3 |
336 |
256 |
322 |
862 |
151 |
606 |
CR |
INV. EX. |
6 |
B |
965 |
399 |
12 |
3 |
335 |
255 |
316 |
882 |
186 |
555 |
CR |
COMP. EX. |
7 |
B |
758 |
63 |
9 |
3 |
337 |
257 |
320 |
914 |
188 |
784 |
CR |
INV. EX. |
8 |
B |
777 |
8 |
6 |
3 |
336 |
256 |
322 |
863 |
199 |
706 |
CR |
COMP. EX. |
9 |
B |
785 |
896 |
11 |
2 |
336 |
256 |
324 |
888 |
193 |
931 |
CR |
INV. EX. |
10 |
B |
783 |
1015 |
11 |
4 |
336 |
256 |
322 |
1000 |
206 |
590 |
CR |
COMP. EX. |
11 |
B |
753 |
311 |
17 |
4 |
336 |
256 |
324 |
960 |
198 |
929 |
CR |
INV. EX. |
12 |
B |
924 |
404 |
25 |
2 |
335 |
255 |
321 |
872 |
212 |
563 |
CR |
COMP. EX. |
13 |
B |
798 |
223 |
12 |
4 |
336 |
256 |
322 |
831 |
162 |
596 |
CR |
INV. EX. |
14 |
B |
792 |
243 |
9 |
2 |
336 |
256 |
327 |
836 |
198 |
954 |
CR |
INV. EX. |
15 |
B |
759 |
207 |
7 |
3 |
335 |
255 |
328 |
878 |
215 |
780 |
CR |
INV. EX. |
16 |
B |
785 |
335 |
11 |
3 |
335 |
255 |
318 |
832 |
203 |
862 |
CR |
INV. EX. |
17 |
B |
779 |
381 |
9 |
4 |
337 |
257 |
327 |
842 |
151 |
693 |
CR |
INV. EX. |
18 |
B |
804 |
203 |
12 |
6 |
337 |
257 |
331 |
865 |
177 |
714 |
CR |
INV. EX. |
19 |
B |
790 |
241 |
7 |
3 |
336 |
256 |
261 |
928 |
156 |
687 |
CR |
INV. EX. |
20 |
B |
803 |
260 |
7 |
4 |
336 |
256 |
20 |
847 |
158 |
671 |
CR |
COMP. EX. |
21 |
B |
761 |
335 |
7 |
3 |
336 |
256 |
409 |
894 |
165 |
571 |
CR |
COMP. EX. |
22 |
B |
797 |
376 |
14 |
4 |
336 |
256 |
631 |
978 |
168 |
603 |
CR |
COMP. EX. |
23 |
B |
771 |
369 |
13 |
3 |
336 |
256 |
316 |
312 |
212 |
510 |
CR |
INV. EX. |
24 |
B |
755 |
222 |
10 |
2 |
336 |
256 |
326 |
284 |
201 |
832 |
CR |
COMP. EX. |
25 |
B |
764 |
414 |
10 |
3 |
336 |
256 |
323 |
34 |
211 |
741 |
CR |
COMP. EX. |
26 |
B |
788 |
404 |
9 |
3 |
337 |
257 |
331 |
915 |
167 |
879 |
CR |
INV. EX. |
27 |
B |
768 |
222 |
11 |
3 |
336 |
256 |
329 |
995 |
111 |
996 |
CR |
INV. EX. |
28 |
B |
768 |
437 |
15 |
3 |
337 |
257 |
325 |
846 |
110 |
639 |
CR |
INV. EX. |
29 |
B |
784 |
321 |
8 |
2 |
337 |
257 |
322 |
910 |
389 |
763 |
CR |
INV. EX. |
30 |
B |
772 |
266 |
13 |
4 |
336 |
256 |
321 |
883 |
398 |
707 |
CR |
INV. EX. |
31 |
B |
755 |
465 |
8 |
3 |
337 |
257 |
328 |
821 |
161 |
23 |
CR |
INV. EX. |
32 |
B |
790 |
311 |
15 |
3 |
336 |
256 |
326 |
897 |
173 |
12 |
CR |
INV. EX. |
33 |
B |
786 |
304 |
8 |
4 |
336 |
256 |
317 |
854 |
210 |
9860 |
CR |
INV. EX. |
34 |
B |
779 |
485 |
7 |
3 |
336 |
256 |
324 |
811 |
196 |
9878 |
CR |
INV. EX. |
35 |
B |
751 |
282 |
7 |
4 |
336 |
256 |
330 |
951 |
196 |
726 |
CR |
INV. EX. |
36 |
B |
808 |
294 |
14 |
3 |
336 |
256 |
328 |
956 |
178 |
828 |
CR |
INV. EX. |
37 |
C |
799 |
211 |
14 |
3 |
336 |
256 |
20 |
899 |
152 |
622 |
CR |
COMP. EX. |
38 |
D |
760 |
208 |
10 |
3 |
294 |
214 |
544 |
875 |
219 |
917 |
CR |
COMP. EX. |
39 |
D |
719 |
490 |
6 |
3 |
295 |
215 |
281 |
818 |
195 |
563 |
CR |
INV. EX. |
Underlines indicate being outside the range of the present invention.
(*)CR: cold rolled steel sheet (no coating), GI: hot-dip galvanized steel sheet (no
alloying of zinc coating), GA: galvannealed steel sheet, EG: electrogalvanized steel
sheet |
[Table 4]
Nos. |
Steels |
Annealing temp. T1 (°C) |
Holding time t1 (s) |
Average cooling rate in temperature range of 750-600°C (°C/s) |
Average cooling rate in temperature range of (Ms+50°C)-quench start temp. T2 (°C/s) |
Ms (°C) |
(Ms-80) (°C) |
Quench start temp. T2 (°C) |
Cooling rate from T2 to 80°C (°C/s) |
Tempering temp. 13 (°C) |
Holding time t3 (s) |
Type* |
|
40 |
D |
935 |
452 |
14 |
3 |
294 |
214 |
276 |
870 |
159 |
565 |
CR |
INV. EX. |
41 |
D |
798 |
77 |
10 |
3 |
294 |
214 |
287 |
938 |
193 |
689 |
CR |
INV. EX. |
42 |
D |
756 |
903 |
15 |
3 |
294 |
214 |
281 |
827 |
171 |
585 |
CR |
INV. EX. |
43 |
D |
805 |
421 |
19 |
3 |
295 |
215 |
278 |
837 |
216 |
737 |
CR |
INV. EX. |
44 |
D |
794 |
453 |
11 |
4 |
294 |
214 |
277 |
876 |
159 |
969 |
CR |
INV. EX. |
45 |
D |
786 |
473 |
6 |
3 |
294 |
214 |
279 |
962 |
151 |
846 |
CR |
INV. EX. |
46 |
D |
770 |
485 |
14 |
4 |
295 |
215 |
285 |
878 |
215 |
607 |
CR |
INV. EX. |
47 |
D |
805 |
453 |
9 |
2 |
294 |
214 |
216 |
877 |
191 |
949 |
CR |
INV. EX. |
48 |
D |
793 |
380 |
12 |
3 |
294 |
214 |
287 |
972 |
187 |
527 |
CR |
INV. EX. |
49 |
D |
751 |
328 |
11 |
4 |
295 |
215 |
282 |
324 |
177 |
723 |
CR |
INV. EX. |
50 |
D |
796 |
292 |
6 |
4 |
294 |
214 |
282 |
822 |
215 |
652 |
CR |
INV. EX. |
51 |
D |
766 |
311 |
10 |
2 |
295 |
215 |
288 |
857 |
114 |
508 |
CR |
INV. EX. |
52 |
D |
808 |
328 |
12 |
3 |
295 |
215 |
281 |
966 |
391 |
882 |
CR |
INV. EX. |
53 |
D |
790 |
421 |
13 |
2 |
295 |
215 |
286 |
890 |
216 |
12 |
CR |
INV. EX. |
54 |
D |
764 |
295 |
11 |
3 |
295 |
215 |
278 |
980 |
198 |
9910 |
CR |
INV. EX. |
55 |
D |
795 |
344 |
13 |
3 |
294 |
214 |
283 |
802 |
177 |
855 |
CR |
INV. EX. |
56 |
D |
804 |
332 |
13 |
3 |
295 |
215 |
284 |
846 |
160 |
813 |
CR |
INV. EX. |
57 |
D |
778 |
361 |
11 |
3 |
295 |
215 |
276 |
836 |
208 |
558 |
CR |
INV. EX. |
58 |
D |
766 |
334 |
8 |
4 |
295 |
215 |
21 |
869 |
169 |
742 |
CR |
COMP. EX. |
59 |
D |
794 |
330 |
7 |
2 |
294 |
214 |
347 |
962 |
188 |
501 |
GA |
COMP. EX. |
60 |
D |
786 |
472 |
13 |
4 |
295 |
215 |
289 |
971 |
168 |
859 |
GA |
INV. EX. |
61 |
D |
756 |
347 |
7 |
4 |
295 |
215 |
277 |
857 |
174 |
747 |
GA |
INV. EX. |
62 |
D |
778 |
220 |
10 |
2 |
295 |
215 |
280 |
845 |
178 |
636 |
EG |
INV. EX. |
63 |
D |
799 |
304 |
7 |
3 |
295 |
215 |
287 |
888 |
199 |
977 |
GA |
INV. EX. |
64 |
D |
801 |
233 |
15 |
2 |
294 |
214 |
288 |
860 |
177 |
518 |
CR |
INV. EX. |
65 |
E |
776 |
328 |
9 |
3 |
298 |
218 |
280 |
905 |
167 |
662 |
CR |
INV. EX. |
66 |
F |
768 |
335 |
10 |
2 |
411 |
331 |
404 |
986 |
167 |
603 |
GA |
INV. EX. |
67 |
G |
766 |
285 |
9 |
3 |
420 |
340 |
413 |
991 |
204 |
546 |
GA |
COMP. EX. |
68 |
H |
775 |
383 |
12 |
3 |
213 |
133 |
199 |
818 |
189 |
812 |
GI |
INV. EX. |
69 |
I |
806 |
292 |
10 |
3 |
183 |
103 |
172 |
915 |
190 |
559 |
GA |
COMP. EX. |
70 |
J |
768 |
465 |
11 |
3 |
334 |
254 |
320 |
970 |
164 |
534 |
GA |
INV. EX. |
71 |
K |
794 |
423 |
8 |
3 |
329 |
249 |
317 |
952 |
212 |
852 |
GA |
COMP. EX. |
72 |
L |
787 |
397 |
10 |
4 |
342 |
262 |
331 |
994 |
174 |
783 |
GA |
INV. EX. |
73 |
M |
773 |
409 |
8 |
4 |
336 |
256 |
325 |
943 |
163 |
646 |
GI |
COMP. EX. |
74 |
N |
785 |
205 |
10 |
3 |
393 |
313 |
383 |
974 |
176 |
537 |
GA |
INV. EX. |
75 |
O |
783 |
203 |
14 |
2 |
408 |
328 |
390 |
896 |
209 |
789 |
GA |
COMP. EX. |
76 |
P |
779 |
359 |
13 |
3 |
272 |
192 |
261 |
813 |
164 |
618 |
GA |
INV. EX. |
77 |
Q |
759 |
297 |
6 |
3 |
260 |
180 |
253 |
957 |
183 |
728 |
GA |
COMP. EX. |
78 |
R |
766 |
394 |
8 |
3 |
329 |
249 |
322 |
879 |
158 |
860 |
GA |
INV. EX. |
Underlines indicate being outside the range of the present invention.
(*)CR: cold rolled steel sheet (no coating), GI: hot-dip galvanized steel sheet (no
alloying of zinc coating), GA: galvannealed steel sheet, EG: electrogalvanized steel
sheet |
[Table 5]
Nos. |
Steels |
Annealing temp. T1 (°C) |
Holding time t1 (s) |
Average cooling rate in temperature range of 750-600°C (°C/s) |
Average cooling rate in temperature range of (Ms+50°C)-quench start temp. T2 (°C/s) |
Ms (°C) |
(Ms-80) (°C) |
Quench start temp. T2 (°C) |
Cooling rate from T2 to 80°C (°C/s) |
Tempering temp. T3 (°C) |
Holding time t3 (s) |
Type* |
|
79 |
S |
784 |
454 |
13 |
3 |
341 |
261 |
327 |
930 |
189 |
933 |
GI |
COMP. EX. |
80 |
T |
801 |
331 |
12 |
2 |
341 |
261 |
325 |
901 |
215 |
887 |
GA |
INV. EX. |
81 |
U |
768 |
416 |
6 |
3 |
347 |
267 |
334 |
919 |
185 |
770 |
GA |
COMP. EX. |
82 |
V |
771 |
480 |
11 |
3 |
336 |
256 |
327 |
824 |
187 |
944 |
GA |
INV. EX. |
83 |
W |
758 |
214 |
12 |
3 |
352 |
272 |
345 |
984 |
164 |
764 |
GA |
COMP. EX. |
84 |
X |
801 |
499 |
13 |
3 |
352 |
272 |
335 |
934 |
171 |
697 |
CR |
INV. EX. |
85 |
Y |
790 |
401 |
6 |
3 |
341 |
261 |
326 |
971 |
194 |
829 |
CR |
COMP. EX. |
86 |
Z |
791 |
235 |
5 |
3 |
339 |
259 |
333 |
890 |
203 |
868 |
GA |
INV. EX. |
87 |
AA |
775 |
413 |
10 |
3 |
329 |
249 |
312 |
865 |
174 |
528 |
GA |
COMP. EX. |
88 |
AB |
787 |
208 |
6 |
3 |
350 |
270 |
333 |
992 |
193 |
882 |
GA |
INV. EX. |
89 |
AC |
798 |
455 |
14 |
3 |
345 |
265 |
340 |
942 |
170 |
856 |
GA |
INV. EX. |
90 |
AD |
777 |
477 |
7 |
4 |
336 |
256 |
330 |
977 |
150 |
508 |
GA |
COMP. EX. |
91 |
AE |
802 |
364 |
11 |
3 |
348 |
266 |
335 |
906 |
156 |
648 |
GA |
INV. EX. |
92 |
AF |
799 |
203 |
12 |
2 |
336 |
258 |
331 |
1000 |
211 |
653 |
GA |
INV. EX. |
93 |
AG |
752 |
276 |
5 |
2 |
326 |
246 |
313 |
979 |
168 |
583 |
CR |
COMP. EX. |
94 |
AH |
804 |
473 |
14 |
3 |
331 |
251 |
320 |
980 |
185 |
523 |
CR |
INV. EX. |
95 |
AI |
786 |
370 |
13 |
2 |
345 |
265 |
336 |
846 |
182 |
795 |
CR |
INV. EX. |
96 |
AJ |
803 |
309 |
10 |
3 |
340 |
260 |
329 |
977 |
220 |
587 |
CR |
COMP. EX. |
97 |
AK |
807 |
472 |
14 |
2 |
340 |
260 |
323 |
900 |
152 |
738 |
CR |
INV. EX. |
98 |
AL |
798 |
340 |
13 |
3 |
340 |
260 |
329 |
834 |
152 |
641 |
CR |
INV. EX. |
99 |
AM |
755 |
292 |
10 |
3 |
325 |
245 |
313 |
843 |
156 |
644 |
CR |
COMP. EX. |
100 |
AN |
705 |
370 |
5 |
3 |
351 |
271 |
340 |
924 |
194 |
645 |
CR |
INV. EX. |
101 |
AO |
927 |
268 |
9 |
4 |
339 |
259 |
320 |
856 |
181 |
676 |
CR |
INV. EX. |
102 |
AP |
773 |
51 |
11 |
2 |
346 |
266 |
338 |
878 |
214 |
580 |
CR |
INV. EX. |
103 |
AQ |
767 |
864 |
7 |
2 |
318 |
236 |
298 |
1000 |
211 |
849 |
CR |
INV. EX. |
104 |
AR |
788 |
208 |
18 |
3 |
326 |
246 |
313 |
942 |
168 |
501 |
CR |
INV. EX. |
105 |
AS |
804 |
235 |
11 |
4 |
340 |
260 |
333 |
948 |
170 |
967 |
CR |
INV. EX. |
106 |
AT |
801 |
287 |
14 |
3 |
321 |
241 |
314 |
888 |
164 |
610 |
CR |
INV. EX. |
107 |
AU |
752 |
236 |
15 |
4 |
349 |
269 |
332 |
963 |
171 |
868 |
CR |
INV. EX. |
108 |
AV |
771 |
260 |
7 |
2 |
333 |
253 |
332 |
807 |
187 |
716 |
CR |
INV. EX. |
109 |
AW |
785 |
472 |
13 |
3 |
332 |
252 |
261 |
921 |
206 |
904 |
CR |
INV. EX. |
110 |
AX |
807 |
354 |
14 |
3 |
345 |
265 |
335 |
325 |
151 |
504 |
CR |
INV. EX. |
111 |
AY |
758 |
450 |
13 |
3 |
333 |
253 |
326 |
833 |
182 |
679 |
CR |
INV. EX. |
112 |
AZ |
803 |
289 |
11 |
3 |
338 |
258 |
321 |
901 |
106 |
597 |
CR |
INV. EX. |
113 |
BA |
790 |
344 |
15 |
2 |
283 |
203 |
274 |
953 |
394 |
563 |
CR |
INV. EX. |
114 |
BB |
782 |
379 |
13 |
3 |
295 |
215 |
277 |
977 |
155 |
23 |
CR |
INV. EX. |
115 |
BC |
797 |
237 |
9 |
4 |
283 |
203 |
273 |
918 |
200 |
9851 |
CR |
INV. EX. |
116 |
BD |
799 |
406 |
11 |
4 |
284 |
204 |
265 |
940 |
196 |
667 |
CR |
INV. EX. |
117 |
BE |
799 |
205 |
8 |
4 |
306 |
226 |
292 |
914 |
168 |
663 |
CR |
INV. EX. |
116 |
BF |
769 |
466 |
8 |
2 |
290 |
210 |
277 |
844 |
213 |
732 |
CR |
INV. EX. |
119 |
BG |
768 |
326 |
12 |
2 |
287 |
207 |
273 |
812 |
153 |
757 |
CR |
INV. EX. |
120 |
BH |
769 |
333 |
7 |
3 |
292 |
212 |
287 |
811 |
150 |
829 |
CR |
INV. EX. |
121 |
BI |
794 |
352 |
13 |
2 |
293 |
213 |
281 |
949 |
207 |
878 |
CR |
INV. EX. |
122 |
BJ |
880 |
310 |
19 |
3 |
331 |
251 |
420 |
1000 |
180 |
800 |
CR |
COMP. EX. |
123 |
BK |
758 |
416 |
8 |
3 |
349 |
269 |
330 |
994 |
166 |
744 |
CR |
INV. EX. |
124 |
BL |
800 |
400 |
14 |
4 |
293 |
213 |
287 |
843 |
181 |
895 |
CR |
INV. EX. |
125 |
BM |
800 |
493 |
6 |
2 |
284 |
204 |
267 |
882 |
207 |
965 |
CR |
INV. EX. |
126 |
BN |
797 |
330 |
6 |
3 |
392 |
312 |
381 |
931 |
209 |
677 |
CR |
INV. EX. |
127 |
BO |
810 |
366 |
8 |
2 |
403 |
323 |
392 |
842 |
181 |
528 |
CR |
INV. EX. |
Underlines indicate being outside the range of the present invention.
(*)CR: cold rolled steel sheet (no coating), GI: hot-dip galvanized steel sheet (no
alloying of zinc coating), GA: galvannealed steel sheet, EG: electrogalvanized steel
sheet |
[0072] The high strength cold rolled steel sheets obtained as described above were used
as test steels. Tensile characteristics, flatness in the width direction, and working
embrittlement resistance were evaluated in accordance with the following test methods.
(Microstructure observation)
[0073] The amount of tempered martensite, the amount of retained austenite, the total amount
of ferrite and bainitic ferrite, and the average grain size of prior austenite were
determined by the methods described hereinabove.
(Proportion of packets having the largest area in prior austenite grains)
[0074] The average proportion of packets having the largest area in prior austenite grains
was determined by the method described hereinabove.
(Tensile test)
[0075] A JIS No. 5 test specimen (gauge length: 50 mm, parallel section width: 25 mm) was
sampled so that the longitudinal direction of the test specimen would be perpendicular
to the rolling direction. A tensile test was performed in accordance with JIS Z 2241
under conditions where the crosshead speed was 1.67 × 10
-1 mm/sec. TS and El were thus measured. In the present invention, 980 MPa or higher
TS was determined to be acceptable.
(Press formability)
[0076] A hole expansion test was performed in accordance with JIS Z 2256 (2010). The steel
sheets obtained were each cut to 100 mm × 100 mm. A 10 mm diameter hole was punched
with a clearance of 12% ± 1%. While holding the steel sheet on a die having an inner
diameter of 75 mm with a blank holder force of 9 tons (88.26 kN), a conical punch
with an apex angle of 60° was pushed into the hole to measure the critical hole diameter
at the occurrence of cracking. The limiting hole expansion ratio λ (%) was determined
from the formula below, and the flangeability was evaluated based on the value of
limiting hole expansion ratio.
Limiting hole expansion ratio: λ (%) = {(Df - D0)/D0} × 100
wherein Df is the hole diameter (mm) at the occurrence of cracking and D0 is the
initial hole diameter (mm).
[0077] Based on the tensile strength (TS), the total elongation (El), and the hole expansion
ratio (λ) obtained as described above, TS × El × λ
0.5/1000 was calculated. The steel sheet was evaluated as "excellent in press formability"
when the calculated value was 80 or more.
(Flatness in the width direction)
[0078] The cold rolled steel sheets obtained as described above were analyzed to measure
the flatness in the width direction. The measurement is illustrated in Fig. 2. Specifically,
a sheet with a length of 500 mm in the rolling direction (coil width × 500 mm L ×
sheet thickness) was cut out from the coil and was placed on a surface plate in such
a manner that the warp at the ends would face upward. The height on the steel sheet
was measured with a contact displacement meter by continuously moving the stylus over
the width. Based on the results, the steepness θ as an index of the flatness of the
steel sheet shape was measured as illustrated in Fig. 2. The flatness was rated as
"×" when the steepness was more than 0.02, as "o" when the steepness was more than
0.01 and 0.02 or less, and as "⊚" when the steepness was 0.01 or less. The steel sheet
was evaluated as "excellent in the flatness in the width direction" when the steepness
was 0.02 or less.
(Working embrittlement resistance)
[0079] The working embrittlement resistance was evaluated by Charpy test. A Charpy test
specimen was a 2 mm deep V-notched test piece that was a stack of steel sheets fastened
together with bolts to eliminate any gaps between the steel sheets. The number of
steel sheets that were stacked was controlled so that the thickness of the stack as
the test piece would be closer to 10 mm. When, for example, the sheet thickness was
1.2 mm, eight sheets were stacked to give a 9.6 mm thick test piece. The sheets for
stacking into the Charpy test specimen were sampled so that the width direction would
be the longitudinal direction. As an index of the working embrittlement resistance,
the ratio vE
0%/vE
10% of the absorbed impact energy at room temperature of the as-produced (unworked) steel
sheet to that of the steel sheet after 10% rolling was measured. The working embrittlement
resistance was rated as "×" when vE
0%/vE
10% was less than 0.6, as "o" when vE
0%/vE
10% was 0.6 or more and less than 0.7, and as "⊚" when vE
0%/vE
10% was 0.7 or more. The Charpy test specimen was evaluated as "excellent in working
embrittlement resistance" when vE
0%/vE
10% was 0.6 or more. Conditions other than those described above conformed to JIS Z 2242:
2018.
[0080] The results are described in Tables 6 to 8. As shown in the tables, INVENTIVE EXAMPLES
achieved 980 MPa or higher TS, excellent press formability, excellent flatness in
the width direction, and excellent working embrittlement resistance. In contrast,
COMPARATIVE EXAMPLES were unsatisfactory in one or more of TS, press formability,
flatness in the width direction, and working embrittlement resistance.
[Table 6]
Nos. |
Steels |
Tempered martensite (area%) |
Retained austenite (vol%) |
Ferrite (area%) |
Bainitic ferrite (area%) |
Total of ferrite and bainitic ferrite (area%) |
Proportion of largest packets in prior austenite grains (area%) |
Prior γ grain size (µm) |
TS (MPa) |
EI (%) |
λ (%) |
TS×EI×λ0.5 /1000 |
Flatness in width direction |
Working embrittlement resistance |
|
1 |
A |
55 |
1 |
36 |
10 |
46 |
50 |
15 |
1602 |
10 |
59 |
123 |
⊚ |
⊚ |
INV. EX. |
2 |
B |
59 |
1 |
33 |
10 |
43 |
53 |
12 |
1791 |
8 |
59 |
110 |
⊚ |
⊚ |
INV. EX. |
3 |
B |
41 |
0 |
49 |
7 |
56 |
47 |
10 |
1006 |
15 |
42 |
98 |
⊚ |
⊚ |
INV. EX. |
4 |
B |
34 |
0 |
60 |
6 |
66 |
50 |
12 |
827 |
18 |
56 |
111 |
⊚ |
⊚ |
COMP. EX. |
5 |
B |
57 |
1 |
46 |
4 |
50 |
60 |
20 |
1758 |
8 |
42 |
91 |
⊚ |
○ |
INV. EX. |
6 |
B |
52 |
0 |
38 |
6 |
44 |
56 |
24 |
1480 |
10 |
58 |
113 |
⊚ |
× |
COMP. EX. |
7 |
B |
43 |
0 |
50 |
7 |
57 |
50 |
12 |
1012 |
15 |
50 |
107 |
⊚ |
⊚ |
INV. EX. |
8 |
B |
35 |
0 |
58 |
8 |
66 |
58 |
14 |
896 |
17 |
60 |
118 |
⊚ |
⊚ |
COMP. EX. |
9 |
B |
51 |
1 |
39 |
6 |
45 |
55 |
18 |
1425 |
11 |
57 |
118 |
⊚ |
○ |
INV. EX. |
10 |
B |
54 |
0 |
41 |
8 |
49 |
55 |
21 |
1540 |
10 |
44 |
102 |
⊚ |
× |
COMP. EX. |
11 |
B |
85 |
0 |
11 |
3 |
14 |
50 |
8 |
1947 |
7 |
40 |
86 |
⊚ |
⊚ |
INV. EX. |
12 |
B |
94 |
1 |
1 |
6 |
7 |
52 |
10 |
2031 |
5 |
50 |
72 |
⊚ |
⊚ |
COMP. EX. |
13 |
B |
51 |
1 |
38 |
9 |
47 |
55 |
10 |
1471 |
10 |
57 |
111 |
⊚ |
⊚ |
INV. EX. |
14 |
B |
57 |
1 |
42 |
7 |
49 |
49 |
10 |
1687 |
9 |
59 |
117 |
⊚ |
⊚ |
INV. EX. |
15 |
B |
59 |
0 |
42 |
8 |
50 |
49 |
14 |
1752 |
8 |
51 |
100 |
⊚ |
⊚ |
INV. EX. |
16 |
B |
54 |
1 |
36 |
9 |
45 |
55 |
13 |
1545 |
10 |
47 |
106 |
⊚ |
⊚ |
INV. EX. |
17 |
B |
84 |
1 |
13 |
7 |
20 |
50 |
14 |
1973 |
6 |
51 |
85 |
⊚ |
⊚ |
INV. EX. |
18 |
B |
92 |
0 |
2 |
2 |
4 |
55 |
9 |
2294 |
4 |
57 |
69 |
⊚ |
⊚ |
INV. EX. |
19 |
B |
49 |
2 |
40 |
7 |
47 |
57 |
14 |
1390 |
11 |
33 |
88 |
⊚ |
⊚ |
INV. EX. |
20 |
B |
52 |
6 |
32 |
10 |
42 |
47 |
10 |
1522 |
10 |
25 |
76 |
⊚ |
⊚ |
COMP. EX. |
21 |
B |
53 |
1 |
39 |
8 |
47 |
95 |
10 |
1557 |
9 |
57 |
106 |
× |
× |
COMP. EX. |
22 |
B |
57 |
0 |
33 |
9 |
42 |
78 |
11 |
1732 |
9 |
49 |
109 |
× |
× |
COMP. EX. |
23 |
B |
56 |
2 |
42 |
6 |
48 |
56 |
12 |
1621 |
9 |
34 |
85 |
⊚ |
⊚ |
INV. EX. |
24 |
B |
52 |
6_ |
40 |
6 |
46 |
59 |
12 |
1458 |
10 |
24 |
71 |
⊚ |
⊚ |
COMP. EX. |
25 |
B |
45 |
7 |
37 |
5 |
42 |
46 |
8 |
1128 |
14 |
21 |
72 |
⊚ |
⊚ |
COMP. EX. |
26 |
B |
52 |
0 |
37 |
7 |
44 |
50 |
12 |
1509 |
10 |
56 |
113 |
⊚ |
⊚ |
INV. EX. |
27 |
B |
55 |
1 |
43 |
5 |
48 |
50 |
11 |
1728 |
9 |
45 |
104 |
⊚ |
○ |
INV. EX. |
28 |
B |
54 |
1 |
34 |
7 |
41 |
48 |
11 |
1684 |
9 |
54 |
111 |
⊚ |
○ |
INV. EX. |
29 |
B |
53 |
0 |
32 |
9 |
41 |
53 |
14 |
1221 |
12 |
43 |
96 |
⊚ |
⊚ |
INV. EX. |
30 |
B |
59 |
0 |
35 |
7 |
42 |
52 |
14 |
1477 |
10 |
42 |
96 |
⊚ |
⊚ |
INV. EX. |
31 |
B |
52 |
0 |
37 |
9 |
46 |
49 |
9 |
1518 |
10 |
46 |
103 |
⊚ |
○ |
INV. EX. |
32 |
B |
52 |
1 |
41 |
7 |
48 |
47 |
10 |
1500 |
10 |
49 |
105 |
⊚ |
○ |
INV. EX. |
33 |
B |
50 |
1 |
37 |
7 |
44 |
46 |
13 |
1354 |
11 |
43 |
98 |
⊚ |
⊚ |
INV. EX. |
34 |
B |
52 |
1 |
37 |
8 |
45 |
58 |
14 |
1465 |
10 |
54 |
108 |
⊚ |
⊚ |
INV. EX. |
35 |
B |
57 |
0 |
36 |
9 |
45 |
56 |
10 |
1690 |
9 |
54 |
112 |
⊚ |
⊚ |
INV. EX. |
36 |
B |
51 |
1 |
39 |
4 |
43 |
53 |
13 |
1447 |
11 |
56 |
119 |
⊚ |
⊚ |
INV. EX. |
37 |
C |
52 |
5 |
37 |
9 |
46 |
50 |
13 |
1480 |
10 |
23 |
71 |
⊚ |
⊚ |
COMP. EX. |
38 |
D |
54 |
1 |
36 |
6 |
42 |
93 |
9 |
1384 |
11 |
48 |
105 |
× |
× |
COMP. EX. |
39 |
D |
41 |
0 |
54 |
5 |
59 |
58 |
14 |
1035 |
15 |
56 |
116 |
⊚ |
⊚ |
INV. EX. |
40 |
D |
54 |
0 |
44 |
4 |
48 |
48 |
18 |
1474 |
10 |
42 |
96 |
⊚ |
○ |
INV. EX. |
41 |
D |
42 |
0 |
54 |
4 |
58 |
56 |
10 |
983 |
15 |
42 |
96 |
⊚ |
⊚ |
INV. EX. |
42 |
D |
58 |
0 |
37 |
4 |
41 |
57 |
16 |
1636 |
9 |
58 |
112 |
⊚ |
○ |
INV. EX. |
43 |
D |
84 |
0 |
11 |
2 |
13 |
50 |
14 |
2138 |
6 |
51 |
92 |
⊚ |
⊚ |
INV. EX. |
44 |
D |
56 |
0 |
38 |
10 |
48 |
56 |
12 |
1564 |
9 |
50 |
100 |
⊚ |
⊚ |
INV. EX. |
45 |
D |
54 |
0 |
41 |
3 |
44 |
56 |
13 |
1486 |
10 |
52 |
107 |
⊚ |
⊚ |
INV. EX. |
46 |
D |
88 |
1 |
14 |
3 |
17 |
59 |
12 |
2120 |
6 |
49 |
89 |
⊚ |
⊚ |
INV. EX. |
47 |
D |
49 |
2 |
42 |
4 |
46 |
59 |
14 |
1201 |
12 |
34 |
84 |
⊚ |
⊚ |
INV. EX. |
Underlines indicate being outside the range of the present invention. |
[Table 7]
Nos. |
Steels |
Tempered martensite (area% ) |
Retained austenite (vol%) |
Ferrite (area%) |
Bainitic ferrite (area%) |
Total of ferrite and bainitic ferrite (area%) |
Proportion of largest packets in prior austenite grains (area%) |
Prior γ grain size (µm) |
TS (MPa) |
EI (%) |
λ (%) |
TSxEIxλ0.5 /1000 |
Flatness in width direction |
Working embrittlement resistance |
|
48 |
D |
52 |
1 |
42 |
4 |
46 |
56 |
14 |
1342 |
11 |
48 |
102 |
⊚ |
⊚ |
INV. EX. |
49 |
D |
57 |
2 |
40 |
8 |
48 |
49 |
14 |
1582 |
9 |
33 |
82 |
⊚ |
⊚ |
INV. EX. |
50 |
D |
55 |
0 |
37 |
8 |
45 |
47 |
11 |
1435 |
10 |
48 |
99 |
⊚ |
⊚ |
INV. EX. |
51 |
D |
55 |
1 |
43 |
3 |
46 |
46 |
9 |
1586 |
10 |
55 |
118 |
⊚ |
○ |
INV. EX. |
52 |
D |
50 |
0 |
37 |
5 |
42 |
52 |
13 |
1046 |
14 |
48 |
101 |
⊚ |
⊚ |
INV. EX. |
53 |
D |
53 |
0 |
35 |
8 |
43 |
52 |
8 |
1343 |
11 |
53 |
108 |
⊚ |
○ |
INV. EX. |
54 |
D |
55 |
1 |
39 |
4 |
43 |
55 |
15 |
1060 |
13 |
52 |
99 |
⊚ |
⊚ |
INV. EX. |
55 |
D |
58 |
1 |
37 |
9 |
46 |
47 |
9 |
1627 |
9 |
42 |
95 |
⊚ |
⊚ |
INV. EX. |
56 |
D |
52 |
0 |
38 |
4 |
42 |
49 |
14 |
1382 |
11 |
55 |
113 |
⊚ |
⊚ |
INV. EX. |
57 |
D |
52 |
1 |
39 |
5 |
44 |
59 |
9 |
1310 |
11 |
52 |
104 |
⊚ |
⊚ |
INV. EX. |
58 |
D |
45 |
5 |
36 |
7 |
43 |
57 |
8 |
1054 |
14 |
26 |
75_ |
⊚ |
⊚ |
COMP. EX. |
59 |
D |
49 |
0 |
39 |
6 |
45 |
88_ |
9 |
1205 |
12 |
50 |
102 |
x |
x |
COMP. EX. |
60 |
D |
52 |
0 |
33 |
10 |
43 |
55 |
14 |
1370 |
11 |
52 |
109 |
⊚ |
⊚ |
INV. EX. |
61 |
D |
57 |
1 |
32 |
9 |
41 |
54 |
13 |
1586 |
10 |
60 |
123 |
⊚ |
⊚ |
INV. EX. |
62 |
D |
51 |
1 |
36 |
7 |
43 |
50 |
15 |
1310 |
12 |
42 |
102 |
⊚ |
⊚ |
INV. EX. |
63 |
D |
51 |
1 |
39 |
6 |
45 |
50 |
13 |
1279 |
12 |
45 |
103 |
⊚ |
⊚ |
INV. EX. |
64 |
D |
52 |
0 |
43 |
3 |
46 |
56 |
15 |
1357 |
11 |
54 |
110 |
⊚ |
⊚ |
INV. EX. |
65 |
E |
59 |
1 |
33 |
7 |
40 |
59 |
10 |
1712 |
9 |
56 |
115 |
⊚ |
⊚ |
INV. EX. |
66 |
F |
42 |
0 |
50 |
5 |
55 |
59 |
10 |
1036 |
14 |
51 |
104 |
⊚ |
⊚ |
INV. EX. |
67 |
G |
33 |
1 |
59 |
5 |
64 |
51 |
11 |
819 |
18 |
57 |
111 |
⊚ |
⊚ |
COMP. EX. |
68 |
H |
53 |
1 |
40 |
8 |
48 |
54 |
11 |
2022 |
7 |
56 |
106 |
⊚ |
○ |
INV. EX. |
69 |
I |
51 |
0 |
46 |
3 |
49 |
55 |
12 |
2035 |
7 |
50 |
101 |
⊚ |
x |
COMP. EX. |
70 |
J |
56 |
0 |
35 |
7 |
42 |
48 |
10 |
1078 |
11 |
52 |
86 |
⊚ |
⊚ |
INV. EX. |
71 |
K |
55 |
1 |
37 |
4 |
41 |
54 |
12 |
820 |
10 |
52 |
59 |
⊚ |
⊚ |
COMP. EX. |
72 |
L |
57 |
2 |
38 |
10 |
48 |
57 |
9 |
1869 |
8 |
33 |
86 |
⊚ |
⊚ |
INV. EX. |
73 |
M |
44 |
6 |
39 |
9 |
48 |
52 |
10 |
1286 |
12 |
26 |
79 |
⊚ |
⊚ |
COMP. EX. |
74 |
N |
42 |
1 |
50 |
10 |
60 |
55 |
12 |
1096 |
13 |
57 |
108 |
⊚ |
⊚ |
INV. EX. |
75 |
O |
35 |
0 |
61 |
4 |
65 |
51 |
10 |
687 |
22 |
49 |
106 |
⊚ |
⊚ |
COMP. EX. |
76 |
P |
59 |
0 |
38 |
4 |
42 |
59 |
15 |
1984 |
8 |
54 |
117 |
⊚ |
○ |
INV. EX. |
77 |
Q |
55 |
1 |
39 |
6 |
45 |
56 |
9 |
1789 |
9 |
54 |
118 |
⊚ |
× |
COMP. EX. |
78 |
R |
52 |
1 |
36 |
6 |
42 |
52 |
9 |
1529 |
10 |
41 |
98 |
⊚ |
○ |
INV. EX. |
79 |
S |
53 |
0 |
38 |
6 |
44 |
48 |
14 |
1533 |
10 |
48 |
106 |
⊚ |
× |
COMP. EX. |
80 |
T |
58 |
0 |
38 |
7 |
45 |
47 |
11 |
1719 |
9 |
47 |
106 |
⊚ |
○ |
INV. EX. |
81 |
U |
57 |
0 |
32 |
9 |
41 |
57 |
12 |
1679 |
9 |
45 |
101 |
⊚ |
× |
COMP. EX. |
82 |
V |
40 |
1 |
48 |
7 |
55 |
54 |
9 |
1066 |
14 |
48 |
103 |
⊚ |
⊚ |
INV. EX. |
83 |
W |
30 |
0 |
64 |
6 |
70 |
45 |
15 |
792 |
19 |
51 |
107 |
⊚ |
⊚ |
COMP. EX. |
84 |
X |
54 |
0 |
40 |
8 |
48 |
51 |
13 |
1586 |
9 |
53 |
104 |
⊚ |
○ |
INV. EX. |
85 |
Y |
51 |
0 |
38 |
6 |
44 |
45 |
14 |
1361 |
11 |
43 |
98 |
⊚ |
× |
COMP. EX. |
86 |
Z |
56 |
1 |
43 |
5 |
48 |
53 |
9 |
1576 |
10 |
48 |
109 |
⊚ |
○ |
INV. EX. |
87 |
AA |
57 |
1 |
37 |
10 |
47 |
46 |
9 |
1713 |
9 |
49 |
108 |
⊚ |
× |
COMP. EX. |
88 |
AB |
53 |
0 |
41 |
4 |
45 |
51 |
14 |
1449 |
10 |
41 |
93 |
⊚ |
⊚ |
INV. EX. |
89 |
AC |
50 |
0 |
44 |
3 |
47 |
55 |
14 |
1496 |
10 |
53 |
109 |
⊚ |
○ |
INV. EX. |
90 |
AD |
58 |
1 |
39 |
9 |
48 |
57 |
14 |
1533 |
10 |
54 |
113 |
⊚ |
× |
COMP. EX. |
91 |
AE |
59 |
1 |
38 |
5 |
43 |
55 |
8 |
1772 |
9 |
45 |
107 |
⊚ |
⊚ |
INV. EX. |
92 |
AF |
54 |
1 |
43 |
3 |
46 |
55 |
13 |
1816 |
8 |
54 |
107 |
⊚ |
○ |
INV. EX. |
93 |
AG |
54 |
0 |
38 |
9 |
47 |
54 |
10 |
1850 |
8 |
59 |
114 |
⊚ |
× |
COMP. EX. |
Underlines indicate being outside the range of the present invention. |
[Table 8]
Nos. |
Steels |
Tempered martensite (area%) |
Retained austenite (vol%) |
Ferrite (area%) |
Bainitic ferrite (area%) |
Total of ferrite and bainitic ferrite (area%) |
Proportion of largest packets in prior austenite grains (area%) |
Prior γ grain size (µm) |
TS (MPa) |
EI (%) |
A (%) |
TS×EI×λ0.5 /1000 |
Flatness in width direction |
Working embrittlement resistance |
|
94 |
AH |
57 |
1 |
37 |
5 |
42 |
51 |
13 |
1656 |
9 |
51 |
106 |
⊚ |
⊚ |
INV. EX. |
95 |
AI |
53 |
0 |
33 |
10 |
43 |
47 |
14 |
1696 |
9 |
49 |
107 |
⊚ |
○ |
INV. EX. |
96 |
AJ |
54 |
1 |
39 |
4 |
43 |
47 |
9 |
1742 |
9 |
56 |
117 |
⊚ |
× |
COMP. EX. |
97 |
AK |
58 |
1 |
44 |
4 |
48 |
48 |
11 |
1739 |
8 |
55 |
103 |
⊚ |
⊚ |
INV. EX. |
98 |
AL |
53 |
0 |
42 |
8 |
50 |
55 |
14 |
1771 |
8 |
57 |
107 |
⊚ |
○ |
INV. EX. |
99 |
AM |
54 |
0 |
40 |
8 |
48 |
57 |
10 |
1810 |
8 |
60 |
112 |
⊚ |
× |
COMP. EX. |
100 |
AN |
40 |
1 |
48 |
7 |
55 |
50 |
9 |
1038 |
14 |
47 |
100 |
⊚ |
⊚ |
INV. EX. |
101 |
AO |
52 |
1 |
41 |
4 |
45 |
55 |
16 |
1466 |
10 |
58 |
112 |
⊚ |
○ |
INV. EX. |
102 |
AP |
43 |
1 |
47 |
9 |
56 |
59 |
8 |
1039 |
15 |
48 |
108 |
⊚ |
⊚ |
INV. EX. |
103 |
AQ |
57 |
1 |
34 |
7 |
41 |
54 |
18 |
1613 |
9 |
43 |
95 |
⊚ |
○ |
INV. EX. |
104 |
AR |
84 |
1 |
11 |
6 |
17 |
49 |
10 |
2081 |
6 |
50 |
88 |
⊚ |
⊚ |
INV. EX. |
105 |
AS |
52 |
1 |
41 |
9 |
50 |
52 |
15 |
1476 |
10 |
56 |
110 |
⊚ |
⊚ |
INV. EX. |
106 |
AT |
58 |
1 |
37 |
4 |
41 |
56 |
15 |
1798 |
9 |
47 |
111 |
⊚ |
⊚ |
INV. EX. |
107 |
AU |
84 |
1 |
14 |
5 |
19 |
48 |
11 |
2071 |
6 |
56 |
93 |
⊚ |
⊚ |
INV. EX. |
108 |
AV |
50 |
2 |
37 |
8 |
45 |
49 |
10 |
1331 |
11 |
45 |
98 |
⊚ |
⊚ |
INV. EX. |
109 |
AW |
55 |
0 |
40 |
6 |
46 |
51 |
8 |
1536 |
10 |
31 |
86 |
⊚ |
⊚ |
INV. EX. |
110 |
AX |
50 |
2 |
40 |
9 |
49 |
47 |
9 |
1444 |
10 |
34 |
84 |
⊚ |
⊚ |
INV. EX. |
111 |
AY |
55 |
1 |
36 |
9 |
45 |
53 |
13 |
1564 |
9 |
53 |
102 |
⊚ |
⊚ |
INV. EX. |
112 |
AZ |
54 |
0 |
42 |
4 |
46 |
47 |
13 |
1682 |
9 |
57 |
114 |
⊚ |
○ |
INV. EX. |
113 |
BA |
50 |
1 |
40 |
7 |
47 |
51 |
9 |
1041 |
15 |
49 |
109 |
⊚ |
⊚ |
INV. EX. |
114 |
BB |
56 |
1 |
39 |
7 |
46 |
47 |
8 |
1603 |
9 |
56 |
108 |
⊚ |
○ |
INV. EX. |
115 |
BC |
52 |
1 |
32 |
9 |
41 |
59 |
15 |
1378 |
11 |
54 |
111 |
⊚ |
⊚ |
INV. EX. |
116 |
BD |
58 |
1 |
33 |
8 |
41 |
53 |
14 |
1633 |
9 |
42 |
95 |
⊚ |
⊚ |
INV. EX. |
117 |
BE |
52 |
0 |
37 |
9 |
46 |
52 |
9 |
1389 |
11 |
52 |
110 |
⊚ |
⊚ |
INV. EX. |
118 |
BF |
51 |
1 |
38 |
6 |
44 |
47 |
12 |
1257 |
12 |
53 |
110 |
⊚ |
⊚ |
INV. EX. |
119 |
BG |
57 |
1 |
39 |
6 |
45 |
50 |
14 |
1632 |
9 |
48 |
102 |
⊚ |
⊚ |
INV. EX. |
120 |
BH |
59 |
1 |
40 |
4 |
44 |
54 |
11 |
1774 |
9 |
40 |
101 |
⊚ |
⊚ |
INV. EX. |
121 |
BI |
54 |
0 |
36 |
4 |
40 |
60 |
8 |
1403 |
11 |
57 |
117 |
⊚ |
⊚ |
INV. EX. |
122 |
BJ |
89 |
0 |
9 |
1 |
10 |
88 |
9 |
1520 |
9 |
45 |
92 |
× |
× |
COMP. EX. |
123 |
BK |
57 |
1 |
37 |
5 |
42 |
45 |
13 |
1743 |
9 |
49 |
110 |
⊚ |
⊚ |
INV. EX. |
124 |
BL |
52 |
1 |
42 |
4 |
46 |
54 |
8 |
1339 |
11 |
47 |
101 |
⊚ |
⊚ |
INV. EX. |
125 |
BM |
51 |
1 |
43 |
4 |
47 |
57 |
8 |
1255 |
12 |
56 |
113 |
⊚ |
⊚ |
INV. EX. |
126 |
BN |
50 |
1 |
36 |
6 |
42 |
58 |
14 |
1406 |
11 |
52 |
112 |
⊚ |
⊚ |
INV. EX. |
127 |
BO |
59 |
0 |
34 |
9 |
43 |
46 |
10 |
1828 |
8 |
51 |
104 |
⊚ |
⊚ |
INV. EX. |
Underlines indicate being outside the range of the present invention. |