Note: Descriptions are shown in the official language in which they were submitted.
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DESCRIPTION
HIGH-STRENGTH STEEL SHEET
Technical Field
The present invention relates to a high-strength steel
sheet having high stretch flangeability after working and
corrosion resistance after painting.
Background Art
Automobile parts, such as chassis and truck frames,
require formability (mainly elongation and stretch
flangeability), and steel having a tensile strength on the
order of 590 MPa has been used for such applications.
However, to reduce the effects of automobiles on the
environment and to improve crashworthiness of automobiles,
use of higher-strength automotive steel sheets has been
promoted in recent years, and use of steel having a tensile
strength on the order of 780 MPa is being investigated.
In general, steel materials having higher strength have
lower workability. High-strength high-workability steel
sheets have therefore been studied. For example, Patent
Documents 1 to 6 describe techniques for improving
elongation and stretch flangeability.
Patent Document 1 discloses a technique relating to
high-workability high-strength steel sheet having a tensile
strength of 590 MPa or more, wherein the steel sheet has a
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substantially ferritic single phase in which carbide
containing Ti and Mo having an average particle size of less
than 10 nm is dispersedly precipitated.
Patent Document 2 discloses a technique relating to a
high-strength hot-rolled steel sheet having a strength of
880 MPa or more and a yield ratio of 0.80 or more. The
steel sheet has a steel structure that contains, on the
basis of mass, C: 0.08% to 0.20%, Si: 0.001% or more but
less than 0.2%, Mn: more than 1.0% but not more than 3.0%,
Al: 0,001% to 0.5%, V: more than 0.1% but not more than 0.5%,
Ti: 0.05% or more but less than 0.2%, and Nb: 0.005% to 0.5%,
provided that the following three formulae are satisfied,
the remainder being Fe and incidental impurities, and that
contains 70% by volume or more ferrite having an average
particle size of 5 m or less and a hardness of 250 Hv or
more.
(Formula 1) (Ti/48 + Nb/93) x C/12 4.5 x 10-5
(Formula 2) 0.5 (V/51 + Ti/48 +
Nb/93)/(C/12) 1.5
(Formula 3) V + Ti x 2 + Nb x 1.4 + C x 2 + Mn x 0.1 >
0.80
Patent Document 3 discloses a technique relating to a
hot-rolled steel sheet that contains, on the basis of mass,
C: 0.05% to 0.2%, Si: 0.001% to 3.0%, Mn: 0.5 to 3.0, P:
0.001% to 0.2%, Al: 0.001% to 3%, and V: more than 0.1% but
not more than 1.5%, the remainder being Fe and impurities,
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and has a structure mainly composed of ferrite phase having
an average particle size in the range of 1 to 5 gm, the
ferrite particles containing carbonitride of V having an
average particle size of 50 nm or less.
Patent Document 4 discloses a thermally stable high-
strength thin steel sheet that contains precipitated carbide
in the steel structure. In the thin steel sheet, carbide
has a NaCl-type crystal structure represented by MC wherein
M denotes a metallic element composed of at least two metals,
and the at least two metals are regularly spaced in a
crystal lattice, forming a superlattice.
Patent Document 5 discloses the following hot-rolled
steel sheet. The steel sheet has a composition of C:
0.0002% to 0.25%, Si: 0.003% to 3.0%, Mn: 0.003% to 3.0%,
and Al: 0.002% to 2.0% on the basis of mass percent, the
remainder being Fe and incidental impurities, the impurities
containing 0.15% or less P, 0.05% or less S, and 0.01% or
less N. A ferrite phase accounts for 70% by area or more of
the metal structure and has an average grain size of 20 gm
or less and an aspect ratio of 3 or less. Seventy percent
or more of ferrite grain boundaries are high-angle grain
boundaries. Among ferrite phases defined by high-angle
grain boundaries, the area percentage of precipitates having
a maximum diameter of 30 gm or less and a minimum diameter
of 5 nm or more is 2% or less of the metal structure.
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Second phases having the largest area percentage among
phases other than the ferrite phases and.the precipitates
have an average grain size of 20 pm or less. High-angle
grain boundaries of ferrite phases are disposed between the
nearest second phases.
Patent Document 6 discloses a drawable high-strength
thin steel sheet that has excellent shape fixability and
burring characteristics, wherein the thin steel sheet
contains, on the basis of mass percent, C: 0.01% to 0.1%, S
0.03%, N 0.005%, and Ti: 0.05% to 0.5%, the Ti content
satisfying Ti-48/12C-48/14N-48/32S 0%,
the remainder being
Fe and incidental impurities, at least the mean values of X-
ray random intensity ratios in a plane at half the thickness
of the steel sheet are 3 or more for {100}<011> to
{223}<110> orientations and 3.5 or less for three
orientations of {554}<225>, {111}<112>, and {111}<110>, the
arithmetical mean roughness Ra of at least one of the
surfaces of the steel sheet ranges from 1 to 3.5, and the
steel sheet is coated with a lubricating composition.
However, the related art described above has the
following problems.
Because the steel sheet contains Mo in Patent Documents
1 and 4, a recent increase in the cost of Mo has resulted in
a marked increase in the cost of the steel sheet.
With the increasing globalization of the automobile
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industry, automotive steel sheets are being used under
severe corrosion conditions, and therefore steel sheets
require higher corrosion resistance after painting. However,
the addition of Mo prevents the formation or growth of
crystals during chemical conversion, thereby lowering the
corrosion resistance of a steel sheet after painting. The
addition of Mo therefore cannot satisfy this requirement.
Thus, the steel described in Patent Documents 1 and 4 does
not have corrosion resistance after painting that satisfies
recent requirements of the automobile industry.
With recent advances in pressing techniques, processing
such as drawing or stretch forming piercing -* flange
forming is increasingly employed. Flanges of steel sheets
formed by such processing require stretch flangeability
after drawing or stretch forming and piercing, that is,
stretch flangeability after working. However, in Patent
Documents 2, 3, and 4, a TS of 780 MPa or more is not always
compatible with sufficient stretch flangeability after
working. The addition of Nb in Patent Document 3
significantly retards the recrystallization of austenite
after hot rolling. Deformed austenite therefore remains in
a steel sheet, thereby lowering workability. The addition
of Nb also disadvantageously increases rolling load in hot
rolling.
Patent Document 5 discloses single-phase ferritic steel
=
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sheets having a tensile strength TS of 422 MPa or less (for
example, test numbers 1 to 5 in Table 6 and test number 45
in Table 8 in Examples) and multiphase steel sheets composed
of a ferrite phase and a second phase and having a tensile
strength TS of 780 MPa or more (for example, test numbers 33
to 36 in Table 6 and test number 49 in Table 8 in Examples).
These steel sheets described in Patent Document 5 mainly
take advantage of solid-solution strengthening due to Si or
Mn and transformation hardening utilizing a hard second
phase. These steel sheets must therefore be cooled to a
temperature in the range of 600 C to 800 C at an average
cooling rate of 30 C/s or more within two seconds after
finish rolling, air-cooled for 3 to 15 seconds, and then
water-cooled at an average cooling rate of 30 C/s or more
before coiling. This promotes two-phase separation during
ferrite transformation, allowing the steel sheets to have a
mixed structure of the ferrite phase and the second phase.
The finish-rolling temperature ranges from (Ae3 point +
10.0 C) to Ae3 point, which is lower than the temperature
range suitable for manufacture according to the present
invention described below. For example, the finish-rolling
temperature for multiphase steel sheets having a tensile
strength TS of 780 MPa or more (test numbers 33 to 36 in
Table 6 in Examples) ranged from 871 C to 800 C. A low
finish-rolling temperature results in a decrease in the
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solubility limit of a carbide-forming element, such as Ti,
in an austenite phase. Furthermore, because rolling
introduces precipitation sites, precipitates having a size
of 20 nm or more are formed. This phenomenon is referred to
as strain-induced precipitation. In the steel sheets and
the method for manufacturing the steel sheets described in
Patent Document 5, strain-induced precipitation increases
the amount of precipitates having a size of 20 nm or more.
Patent Document 5 also discloses a technique in which a
ferritic single phase can be manufactured by greatly
decreasing the C content and decreasing the amount of
austenite forming element, Mn, in a steel composition (see
steel numbers AA to AE in Table 2 in Examples). However, a
decrease in the amount of Mn, which is also a solid-solution
strengthening element, lowers the solid-solution
strengthening level. A decrease in C content results in a
decrease in the amount of precipitated carbide, for example,
of Ti or Nb, which has precipitation hardening effects,
thereby lowering the precipitation hardening level. Thus,
even with a combination of the solid-solution strengthening
level and the precipitation hardening level, a single-phase
ferritic steel sheet cannot have a strength of 780 MPa or
more (see test numbers 1 to 5 in Table 6 and test number 45
in Table 8 in Examples). For these reasons, an object of
the present invention, that is, a steel sheet that has a
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substantially ferritic single phase, a tensile strength of
780 MPa or more, and other characteristics cannot be
manufactured by the technique described in Patent Document 5.
Patent Document 6 discloses steel sheets having a
tensile strength aB of 780 MPa or more (for example, steel
symbols A-4, A-8, A-10, C, E, and H in Table 2 in Examples).
The YRs of these steel sheets (YR represents ay/as x 100 (%))
are as low as 69% to 74%, indicating that these steel sheets
contain a hard second phase, such as a martensite phase.
As in Patent Document 5, the possible basic ideas
behind the design of a steel sheet having a strength of 780
MPa or more according to Patent Document 6 mainly take
advantage of solid-solution strengthening due to Si or Mn
and transformation hardening utilizing a hard second phase.
As described in Patent Document 5, therefore, rolling at a
total reduction of 25% or more must be performed at a
finish-rolling temperature (Ar3 point + 100 C or less) lower
than the temperature range suitable for manufacture
according to the present invention described below. For
example, according to an example of Patent Document 6, the
finish-rolling temperature for a steel sheet having a
tensile strength aB of 780 MPa or more ranged from 800 C to
890 C. In the steel sheets and the method for manufacturing
the steel sheets described in Patent Document 6, as
described in Patent Document 5, strain-induced precipitation
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increases the amount of precipitates having a size of 20 nm
or more. Consequently, an object of the present invention,
that is, a steel sheet that has a substantially ferritic
single phase, a tensile strength of 780 MPa or more, and
other characteristics cannot be manufactured.
Patent Document 1: Japanese Patent No. 3591502
Patent Document 2: Japanese Unexamined Patent
Application Publication No. 2006-161112
Patent Document 3: Japanese Unexamined Patent
Application Publication No. 2004-143518
Patent Document 4: Japanese Unexamined Patent
Application Publication No. 2003-321740
Patent Document 5: Japanese Unexamined Patent
Application Publication No. 2003-293083
Patent Document 6: Japanese Unexamined Patent
Application Publication No. 2003-160836
Disclosure of Invention
In view of the situations described above, it is an
object of the present invention to provide a high-strength
steel sheet having high stretch flangeability after working
and corrosion resistance after painting.
As a result of investigations to develop a high-
strength hot-rolled steel sheet that has high stretch
flangeability after working, corrosion resistance after
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painting, and a tensile strength of 780 MPa or more, the
present inventors obtain the following findings.
i) To manufacture a high-strength steel sheet having
high corrosion resistance after painting, precipitates must
remain fine (less than 20 nm), and the percentage of fine
precipitates (having a size less than 20 nm) must be
increased. Although precipitates containing Ti-Mo or Ti-V
remain fine, mixed precipitation of Ti and V is useful in
improving corrosion resistance after paining.
ti) Solid solution of V is important in improving
stretch flangeability after working. There is an optimum V
content of solid solution for an improvement in
characteristics.
The present invention has been accomplished on the
basis of these findings and is summarized as follows:
[1] A high-strength steel sheet comprising, on the
basis of mass percent, C: 0.02% to 0.20%, Si: 0.3% or less,
Mn: 0.5?s to 2.5%, P: 0.06% or less, S: 0.01% or less, Al:
0.14, or less, Ti: 0.05% to 0.25%, and V: 0.05% to 0.25%,
the remainder being Fe and incidental impurities, wherein
the steel sheet has a ferritic single phase, the ferritic
single phase containing precipitates having a size of less
than 20 nm, the precipitates containing 200 to 1750 mass
ppm Ti and 150 to 1750 mass ppm V, V dissolved in solid
solution being 200 or more but less than 1750 mass ppm, the
ratio of the Ti content to the V content of precipitates
having a size of less than 20 nm satisfies 0.4 (Ti/48) /
(V/51) 2.5, and tensile strength, TS, is not less than
780 MPa.
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[2] A high-strength steel sheet comprising, on the
basis of mass percent, C: 0.02% to 0.20%, Si: 0.3% or less,
Mn: 0.5% to 2.5%, P: 0.06% or less, S: 0.01% or less, Al:
0.1.6 or less, Ti: 0.05% to 0.25%, V: 0.05% to 0.25%, and
any one or two or more of Cr: 0.01% to 0.5%, W: 0.005% to
0.2%, and Zr: 0.0005% to 0.05%, the remainder being Fe and
incidental impurities, wherein the structure of the steel
sheet has a ferritic single phase, the ferritic single
phase containing precipitates having a size of less than 20
nm, the precipitates containing 200 to 1750 mass ppm Ti and
150 to 1750 mass ppm V, V dissolved in solid solution being
200 or more but less than 1750 mass ppm and tensile
strength, TS, is not less than 780 MPa.
F3] The high-strength steel sheet according to [1] or
[2], wherein the steel sheet has a one-side maximum peel
width of 3.0 mm or less after a tape peel test in a warm
salt water immersion test.
[4] The high-strength steel sheet according to [1] or
[2], wherein the steel sheet has a stretch flangeability An
of 60;:, or more after rolling at an elongation percentage of
10%.
In the present specification, the percentages and ppm
of components of steel are based on mass percent and mass
ppm. High-strength steel sheets according to the present
invention have a tensile strength (hereinafter also
referred to as TS) of 780 MPa or more and include hot-
rolled steel sheets and surface-treated steel sheets, which
are high-strength steel sheets subjected to surface
treatment, such as plating.
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Target characteristics of the present invention include
a stretch flangeability (X10) of 60% or more after rolling at
an elongation percentage of 10% and a one-side maximum peel
width of 3.0 mm or less after a tape peel test in a warm
salt water immersion test (SDT) described below.
The present invention provides a high-strength hot-
rolled steel sheet that has high stretch flangeability after
working, corrosion resistance after painting, and a TS of
780 MPa or more. The present invention has these advantages
without the addition of Mo and can therefore reduce costs.
For example, use of a high-strength hot-rolled steel
sheet according to the present invention in automobile
chassis and truck frames should allow thickness reduction,
reduce the effects of automobiles on the environment, and
markedly improve crashworthiness of automobiles.
Best Modes for Carrying Out the Invention
The present invention will be described in detail below.
(1) First, the reason to limit the chemical components
(composition) of steel according to the present invention
will be described below.
C: 0.02% to 0.20%
C can be precipitated in ferrite as carbide with Ti or
V, thereby contributing to high strength of a steel sheet.
0.02% or more C is required to achieve a TS of 780 MPa or
more. However, more than 0.20% C results in coarsening of
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precipitates and the formation of a second phase, lowering
stretch flangeability after working. Thus, the C content
ranges from 0.02% to 0.20%, preferably 0.03% to 0.15%.
Si: 0.3% or less
Although Si can contribute to solid-solution
strengthening, the addition of more than 0.3% Si results in
the formation of cementite at grain boundaries, lowering
stretch flangeability after working. Thus, the Si content
is 0.3% or less, preferably 0.001% to 0.2%.
Mn: 0.5% to 2.5%
Mn can contribute to solid-solution strengthening.
However, the TS is less than 780 MPa at a Mn content of less
than 0.5%. The addition of more than 2.5% Mn markedly
lowers weldability. Thus, the Mn content ranges from 0.5%
to 2.5%, preferably 0:6% to 2.0%.
P: 0.06% or less
P can segregate at prior austenite grain boundaries,
lowering workability and low-temperature toughness. Thus,
the P content is preferably minimized and is 0.06% or less,
preferably in the range of 0.001% to 0.055%.
S: 0.01% or less
S can segregate at prior austenite grain boundaries or
can be precipitated as MnS. The segregation or a large
amount of MnS lowers low-temperature toughness. S also
markedly lowers stretch flangeability, regardless of the
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presence or absence of working. Thus, the S content is
preferably minimized and is 0.01% or less, preferably in the
range of 0.0001% to 0.005%.
Al: 0.1% or less
Al can be added to steel as a deoxidizer and
effectively improves the cleanliness of the steel.
Preferably, 0.001% or more Al is added to steel to produce
this effect. However, more than 0.1% Al results in the
generation of a large number of inclusions, causing flaws in
a steel sheet. Thus, the Al content is 0.1% or less,
preferably 0.01% to 0.04%.
Ti: 0.05% to 0.25%
Ti is very important for the precipitation hardening of
ferrite and is an important factor for the advantages of the
present invention. A required strength is difficult to
achieve at a Ti content of less than 0.05%. However, the
effects of Ti become saturated at a Ti content of more than
0.25%, and more than 0.25% Ti only increases costs. Thus,
the Ti content ranges from 0.05% to 0.25%, preferably 0.08%
to 0.20%.
V: 0.05% to 0.25%
V can contribute to an improvement in strength by
precipitation hardening or solid-solution strengthening.
Like Ti, V is therefore an important factor for the
advantages of the present invention. A proper amount of V,
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together with Ti, tends to be precipitated as fine Ti-V
carbide having a partiele size (hereinafter also referred to
as "size") of less than 20 nm. Unlike Mo, V does not lower
corrosion resistance after painting. Less than 0.05% V is
insufficient for the effects described above. However, the
effects of V become saturated at a V content of more than
0.25%, and more than 0.25% V only increases costs. Thus,
the V content ranges from 0.05% to 0.25%, preferably 0.06%
to 0.20%.
With these essential additive elements, steel according
to the present invention can have target characteristics.
In addition to the essential additive elements, any one or
two or more of Cr: 0.01% to 0.5%, W: 0.005% to 0.2%, and Zr:
0.0005% to 0.05% may be added for the following reasons.
Cr: 0.01% to 0.5%, W: 0.005% to 0.2%, and Zr: 0.0005%
to 0.05%
Like V, Cr, W, and Zr can strengthen ferrite as a
precipitate or solid solution. Less than 0.01% Cr, less
than 0.005% W, or less than 0.0005% Zr makes a negligible
contribution to high strength of steel. However, more than
0.5% Cr, more than 0.2% W, or more than 0.05% Zr lowers
workability. Thus, when any one or two or more of Cr, W,
and Zr are added, their amounts are Cr: 0.01% to 0.5%, W:
0.005% to 0.2%, and Zr: 0.0005% to 0.05%, preferably Cr:
0.03% to 0.3%, W: 0.01% to 0.18%, and Zr: 0.001% to 0.04%.
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The remainder consists of Fe and incidental impurities.
As an incidental impurity, for example, 0 forms a non-
metallic inclusion and has adverse effects on the quality of
steel. 0 is therefore desirably decreased to 0.003% or less.
In the present invention, 0.1% or less Cu, Ni, Sn, and/or Sb
may be contained as a trace element without compromising the
operational advantages of the present invention.
(2) The structure of a high-strength steel sheet
according to the present invention will be described below.
Substantially Ferritic Single Phase =
To achieve a TS of 780 MPa or more and improve stretch
flangeability after working, ferrite having a low
dislocation density is effective, and a single phase is
effective. In particular, a highly ductile ferritic single
phase has a marked improving effect on stretch flangeability
after working. However, a completely ferritic single phase
is not necessary, and even a substantially ferritic single
phase can sufficiently produce the effect. A substantially
ferritic single phase, as used herein, refers to allowance
for a minute amount of another phase or precipitate other
than carbide of the present invention, and the volume
percentage of ferrite is preferably 95% or more. A
substantially ferritic single phase may contain up to 5% by
volume of cementite, pearlite, and/or bainite without
affecting the characteristics of the present invention.
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The volume percentage of ferrite can be determined by
exposing a microstructure in the vertical cross-section
parallel to the rolling direction using 3% nital, observing
the microstructure at a quarter thickness in the depth
direction with a scanning electron microscope (SEM) at a
magnification of 1500, and determining the ferrite area
ratio, for example, using an image-processing software
"Ryusi Kaiseki (particle analysis) II" from Sumitomo Metal
Technology, Inc.
200 to 1750 ppm Ti and 150 to 1750 ppm V in
Precipitates Having a Size below 20 nm in a Ferritic Single
Phase
In a high-strength steel sheet according to the present
invention, precipitates containing Ti and/or V exist in
ferrite mainly as carbides. This is probably because the
solubility limit of C in ferrite is low, and supersaturated
C is therefore easily precipitated in ferrite as carbide.
Such a precipitate increases the hardness (strength) of soft
ferrite, thereby achieving a TS of 780 MPa or more. Such a
precipitate also increases YS, achieving YR (= YS/YR) of 83%
or more.
As described above, to manufacture a high-strength
steel sheet, it is important that precipitates remain fine
(less than 20 nm), and the percentage of fine precipitates
(having a size less than 20 nm) is increased. A precipitate
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having a size of 20 nm or more has a small effect in
preventing dislocation movement and cannot sufficiently
increase the hardness of ferrite, sometimes resulting in low
strength.
A further investigation revealed that a fine
precipitate size is important for corrosion resistance after
painting. In conventional Ti (addition of Ti alone) HSLA
steel, a precipitate have a tendency to become coarse with
increasing Ti content. In such a steel sheet, therefore,
corrosion resistance after painting also has a tendency to
decrease with decreasing strength. Although the reason for
a deterioration in corrosion resistance after painting
associated with coarsening of a precipitate is not clear, a
coarse precipitate should prevent the formation or growth of
crystals during chemical conversion.
Thus, a precipitate preferably has a size of less than
20 nm. A fine precipitate having a size of less than 20 nm
can be formed by the addition of both Ti and V. V forms a
complex carbide .mainly with Ti. Although there is no clear
reason, these precipitates remain stable and fine at high
temperatures within the coiling temperature within the scope
of the present invention for a long period of time.
It is important to control the Ti content and the V
content of precipitates having a size of less than 20 nm.
When the Ti content and the V content of precipitates having
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a size of less than 20 nm are less than 200 ppm and less
than 150 ppm, respectively, the number density of the
precipitates is small, and the distance between precipitates
increases. The precipitates therefore have a small effect
in preventing dislocation movement. Thus, the precipitates
cannot sufficiently increase the hardness of ferrite, and
therefore the TS cannot be 780 MPa or more. When the Ti
content and the V content of precipitates having a size of
less than 20 nm are 200 ppm or more and less than 150 ppm,
respectively, the precipitates have a tendency to become
coarse, and therefore the TS may be less than 780 MPa. When
the Ti content and the V content of precipitates having a
size of less than 20 nm are less than 200 ppm and 150 ppm or
more, respectively, the precipitation efficiency of V
decreases, and therefore the TS may be less than 780 MPa.
When the Ti content or the V content of precipitates having
a size of less than 20 nm is more than 1750 ppm, the
corrosion resistance after painting decreases, and therefore
the target characteristics cannot be achieved. This is
probably because a large number of fine precipitates prevent
the formation or growth of crystals on the surface of a
steel sheet during chemical conversion. Thus, the amounts
of precipitated Ti and V in precipitates having a size of
less than 20 nm must be satisfactorily controlled.
When the ratio of the Ti content to the V content of
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precipitates having a size of less than 20 nm satisfies 0.4
(Ti/48)/(V/51) 2.5, the TS can be 785 MPa or more, thus
achieving more suitable, conditions. Although there is no
clear reason, optimization of the ratio of Ti to V should
improve heat stability.
Thus, the Ti content and the V content of precipitates
having a size of less than 20 nm range from 200 to 1750 ppm
and 150 to 1750 ppm, respectively. Furthermore, the ratio
of the Ti content to the V content of precipitates having a
size of less than 20 nm preferably satisfies 0.4
(Ti/48)/(V/51) 2.5.
A precipitate and/or an inclusion is hereinafter also
collectively referred to as a precipitate or the like.
The Ti content and the V content can be controlled by
the coiling temperature. The coiling temperature preferably
ranges from 500 C to 700 C. At a coiling temperature above
700 C, precipitates become coarse, and the amounts of
precipitated Ti and V in precipitates having a size of less
than 20 nm are less than 200 ppm and less than 150 ppm,
respectively, and the TS cannot be 780 MPa or more. At a
coiling temperature below 500 C, the amounts of precipitated
Ti and V in precipitates having a size of less than 20 nm
are also less than 200 ppm and less than 150 ppm,
respectively. Such a low coiling temperature should result
in insufficient diffusion of Ti and V.
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The Ti content and the V content of precipitates having
a size of less than 20 nm can be determined by the following
method.
After a predetermined amount of sample is electrolyzed
in an electrolyte, the sample is removed from the
electrolyte and is immersed in a dispersive solution.
Precipitates in the solution is filtered with a filter
having a pore size of 20 nm. Precipitates in filtrate
passing through the filter having a pore size of 20 nm have
a size of less than 20 nm. The filtrate after filtration is
appropriately analyzed by inductively coupled plasma (ICP)
emission spectroscopic analysis, ICP mass spectrometry,
atomic absorption spectrometry, or the like to determine the
Ti content and the V content of precipitates having a size
of less than 20 nm.
Structure Containing 200 ppm or More but Less Than 1750
ppm V in Solid Solution
In the present invention, V in solid solution is the
most important factor. Solid solution of V is important in
improving stretch flangeability after working. Less than
200 ppm V in solid solution has an insufficient effect, and
200 ppm or more V in solid solution is required to-produce
the effect described above. 1750 ppm or more V in solid
solution exhibits a saturated effect and is considered as an
upper limit.
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Thus, the amount of V in solid solution is 200 ppm or
more but less than 1750 ppm. Although the workability of
steel according to the present invention slightly
deteriorates with increasing strength, when the Ti content
and the V content of precipitates having a size of less than
20 nm are both 1750 ppm or less, 200 ppm or more V in solid
solution can sufficiently ensure.target stretch
flangeability after working.
200 ppm or more but less than 1750 ppm V in solid
solution can be measured, for example, by the following
method.
After a predetermined amount of sample is electrolyzed
in a nonaqueous solvent electrolyte, the electrolyte is
subjected to elementary analysis. The analysis method may
be inductively coupled plasma (ICP) emission spectroscopic
analysis, ICP mass spectrometry, or atomic absorption
spectrometry.
(3) A method for manufacturing a high-strength steel
sheet according to the present invention will be described
below.
For example, a high-strength steel sheet according to
the present invention can be manufactured by heating a steel
slab adjusted within the chemical component ranges described
above at a temperature in the range of 1150 C to 1350 C, hot-
rolling the steel slab at a finish-rolling temperature in
CA 02693489 2010-01-15
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the range of 850 C to 1100 C, and coiling the rolled steel at
a temperature in the range of 500 C to 700 C. Conditions
suitable for these processes will be described in detail
below.
Steel Slab Heating Temperature: 1150 C to 1350 C
A carbide-forming element, such as Ti or V, is mostly
present as a precipitate in a steel slab. To be
precipitated as desired in a ferrite phase after hot rolling,
a precipitate in the form of carbide must be temporarily
dissolved before hot rolling. A precipitate must therefore
be heated at 1150 C or more.
At a temperature below 1150 C, carbide having a size of
20 nm or more, which does not contribute to precipitation
hardening or corrosion resistance after painting, remains.
This reduces the amount of Ti and V involved in the
formation of fine precipitates having a size of less than 20
nm required for the advantages of the present invention. A
target amount of precipitates having a size of less than 20
nm cannot therefore be obtained in coiling described below.
In a method for manufacturing a steel sheet according to the
present invention, most desirably, carbide containing Ti or
V remains dissolved during slab heating and finish rolling,
and is precipitated as fine carbide containing Ti or V
during coiling after finish rolling. The heating
temperature is therefore more preferably 1170 C or more so
CA 02693489 2010-01-15
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that carbide can be dissolved almost completely.
However, heating at a temperature above 1350 C
excessively increases the crystal grain size, lowering
stretch flangeability and elongation after working. Taking
subsequent heat-treatment conditions into consideration, an
increase in crystal grain size can be almost completely
prevented at a heating temperature of 1300 C or less.
Thus, the slab heating temperature preferably ranges
from 1150 C to 1350 C, more preferably 1170 C to 1300 C.
Finish-Rolling Temperature in Hot Rolling: 850 C to
1100 C
The control of finish-rolling temperature is important
in ensuring the Ti content and the V content of precipitates
having a size of less than 20 nm according to the present
invention. Preferably, a steel slab after working is hot-
rolled at a finish-rolling temperature in the range of 850 C
to 1100 C, which is the final temperature of hot rolling. At
a finish-rolling temperature below 850 C, a steel slab is
rolled in a ferrite + austenite region and has an elongated
ferrite phase. This may lower stretch flangeability or
elongation after working. Even if a steel slab is heated at
a temperature of 1150 C or more to temporarily dissolve a
carbide precipitate before rolling, carbide containing Ti or
V is precipitated at a finish-rolling temperature below
850 C because of strain-induced precipitation. This reduces
=
CA 02693489 2010-01-15
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the amount of Ti and V involved in the formation of fine
precipitates having a size of less than 20 nm required for
the advantages of the present invention. A target amount of
precipitates having a size of less than 20 nm cannot
therefore be obtained in coiling described below. Thus, it
is important to perform the subsequent coiling process while
carbide containing Ti or V temporarily dissolved during the
slab heating described above remains dissolved in finish
rolling as much as possible. The finish-rolling temperature
is more preferably 935 C or more such that carbide remains
dissolved.
A finish-rolling temperature above 1100 C may result in
coarsening of ferrite particles and a TS below 780 MPa. The
finish-rolling temperature is more preferably 990 C or less
to prevent coarsening of ferrite particles.
Thus, the finish-rolling temperature preferably ranges
from 850 C to 1100 C, more preferably 935 C to 990 C.
Coiling Temperature: 500 C to 700 C
The control of coiling temperature is important in
ensuring the Ti content and the V content of precipitates
having a size of less than 20 nm in the present invention.
As described above, this is because, in the most desirable
manufacturing form, this coiling process yields a large
number of precipitation sites from which carbide is
precipitated, thus preventing carbide grains from growing to
CA 02693489 2010-01-15
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20 nm or more. The coiling temperature preferably ranges
from 500 C to 700 C so that steel has a substantially
ferritic single phase and the characteristics of the present
invention can be achieved.
In the present invention, a coiling temperature below
500 C may result in an insufficient amount of precipitated
carbide containing Ti and/or V and reduced strength.
Furthermore, a bainite phase may be formed in place of a
ferritic single phase.
To form a large number of precipitation sites and
produce carbide from these precipitation sites, the coiling
temperature is preferably 500 C or more, more preferably
550 C or more.
A coiling temperature above 700 C may result in
coarsening of precipitated carbide and reduced strength. A
coiling temperature above 700 C may also promote the
formation of a pearlite phase, lowering stretch
flangeability after working. The coiling temperature is
more preferably 650 C or less to prevent coarsening of
precipitated carbide without fail.
Thus, the coiling temperature preferably ranges from
500 C to 700 C, more preferably 550 C to 650 C.
Steel sheets according to the present invention include
surface-treated steel sheets and surface-coated steel sheets.
In particular, a steel sheet according to the present
CA 02693489 2010-01-15
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invention may be subjected to hot-dip galvanizing to form a
galvanized steel sheet, and the present invention can be
suitably applied to such a galvanized steel sheet. Because
a steel sheet according to the present invention has
excellent workability, such a galvanized steel sheet can
also have excellent workability. Hot-dip galvanizing is
zinc and zinc-based (approximately 90% or more) hot dipping
and includes hot dipping including an alloying element, such
as Al or Cr, as well as zinc. Hot-dip galvanizing may be
performed alone or followed by alloying.
A steel melting method is not particularly limited, and
any known melting method may be suitable. For example, a
suitable melting method involves melting in a converter or
an electric furnace and secondary refining in a vacuum
degassing furnace. A casting method is preferably
continuous casting in terms of productivity and quality.
After casting, hot direct rolling may be performed
immediately or after concurrent heating, without
compromising the advantages of the present invention.
Furthermore, a hot-rolled material may be heated after rough
rolling and before finish rolling, continuous hot rolling in
which rolled materials are joined may be performed after
rough rolling, or heating and continuous rolling of a
heating material of a rolled material may be performed
simultaneously. These do not compromise the advantages of
CA 02693489 2010-01-15
28 -
the present invention.
EXAMPLES
EXAMPLE 1
Steel having a composition shown in Table I was melted
in a converter and was formed into a steel slab by
continuous casting. The steel slab was subjected to heating,
hot rolling, and coiling under conditions shown in Table 2
to form a hot-rolled steel sheet having a thickness of 2.0
mm.
Table 1
Type Composition (mass%)
ofNote
Si Mn P S Al Ti V
steel
A 0.040 0.01 1.45 0.01 0.0015 0.03 0.105 0.120
Conforming steel
B 0.120
0.02 1.20 0.02 0.0008 0.03 0.240 0.100 Conforming steel
C 0.100 0.02 1.20 0.01 0.0080 0.03 0.110 0.245
Conforming steel
D 0.150
0.02 1.40 0,03 0.0020 0.03 0.230 0.224 Conforming steel
E 0.050
0.01 2.02 0.01 0.0020 0.03 0.120 0.120 Conforming steel
F 0.050
0.01 0.65 0.01 0.0015 0.03 0.110 0.136 Conforming steel
G 0.045
0.02 1.34 0.02 0.0007 0.02 0.060 0.110 Conforming steel
H 0.050
0.02 1.30 0.01 0.0008 0.02 0.110 0.052 Conforming steel
0
I 0.030 0.01 1.32 0.01 0.0007 0.02 0.080 0.070
Conforming steel
J 0.040
0.01 1.40 0.02 0.0015 0.03 0.126 0.152 Conforming steel
K 0.250
0.01 1.20 0.02 0.0020 0.03 0.120 0.130 Nonconforming co
ks)
L 0.001 0.01 1.19 0.02 0.0020 0.03 0.120 0.130 ,Nonconforming
0
M 0.080 0.50 1.30 0.01 0.0012 0.03 0.070 0.070 Nonconforming 0
0
N 0.050 0.01 0.35 0.02 0.0015 0.03 0.080 0.080 Nonconforming
O 0.050
0.01 3.00 0.02 0.0014 0.03 0.080 0.080 Nonconforming
P 0.150
0.01 1.60 0.02 0.0015 0.03 0.040 0.120 Nonconforming
Q 0.160 0.01 1.60 0.02 0.0016 0.02 0.070 0.032 Nonconforming
R 0.152 0.01 1.62 0.02 0.0015 0.03 0.280 0.120 Nonconforming
S 0.161
0.01 1.61 0.02 0.0014 0.03 0.150 0.300 Nonconforming
X 0.090 0.06 1.35 0.04 0.0014 0.05 0.150 0.160
Conforming steel
CA 02693489 2010-01-15
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The microstructure of the hot-rolled steel sheet was
analyzed by the following method to determine the Ti content
and the V content of precipitates having a size of less than
20 nm and the amount of V in solid solution. The tensile
strength TS, the stretch flangeability after working
and
the corrosion resistance after painting (SDT one-side kio,
maximum peel width) were measured.
Analysis of Microstructure
The hot-rolled steel sheet thus formed was cut into an
appropriate size. Approximately 0.2 g of hot-rolled steel
sheet was subjected to constant-current electrolysis at an
electric current density of 20 mA/cm2 in 10% AA electrolyte
(10% by volume acetylacetone-1% by mass tetramethylammonium
chloride-methanol).
Measurement of the Ti Content and the V Content of
Precipitates Having a Size of Less Than 20 nm
After electrolysis, a test piece on which a precipitate
was deposited was removed from the electrolyte and was
immersed in aqueous sodium hexametaphosphate (500 mg/1)
(hereinafter referred to as aqueous SHMP). Ultrasonic
vibration was applied to the test piece to detach and
extract the precipitate from the test piece in aqueous SHMP.
The aqueous SHMP containing the precipitate was then passed
through a filter having a pore size of 20 nm. The filtrate
was analyzed with an ICP spectrometer to measure the
CA 02693489 2010-01-15
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absolute amounts of Ti and V in the filtrate. The absolute
amounts of Ti and V were divided by the weight of the
electrolyzed sample to calculate the Ti content and the V
content of precipitates having a size of less than 20 nm.
The weight of electrolyzed sample was calculated by
subtracting the sample weight after the detachment of the
precipitate from the sample weight before electrolysis.
Measurement of the Amount of V in Solid Solution
After electrolysis, the concentrations of V and a
comparative element Fe in the electrolyte were measured by
ICP mass spectrometry. On the basis of the concentrations
thus measured, the ratio of the concentration of V to the
concentration of Fe was calculated. The ratio was
multiplied by the Fe content of the sample to calculate the
amount of V in solid solution. The Fe content of the sample
can be calculated by subtracting the summation of
compositions other than Fe from 100%.
TS
A tensile test according to JIS Z 2241 was performed
with a JIS No. 5 specimen in the tensile direction parallel
to the rolling direction to measure TS.
Stretch flangeability after Working: Xn
After rolling at an elongation percentage of 10%, a
hole expanding test according to the Japan Iron and Steel
Federation Standard JFS T 1001 was performed to measure Xn.
Mk 02693489 2012-02-09
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Corrosion Resistance after Painting: SDT One-Side
Maximum Peel Width
A chemical conversion treatment was performed under
more adverse temperature and concentration conditions than
the standard conditions using a degreasing agent,
SurfcleanerTM EC090, a surface conditioner, SurffineTM 5N-10,
and a chemical conversion treatment agent, SurfdineTM
SD2800, all manufactured by Nippon Pint Co., Ltd. As an
example of standard conditions, a degreasing process
included a concentration of 16 g/l, a treatment temperature
in the range of 42 C to 44 C, a treatment time of 120 s,
and spray degreasing, and a surface conditioning process
included a total alkalinity in the range of 1.5 to 2.5
points, a free acidity in the range of 0.7 to 0.9 points,
an accelerator concentration in the range of 2.8 to 3.5
points, a treatment temperature of 44 C, and a treatment
time of 120 s. Under adverse conditions, a treatment
temperature in a chemical conversion treatment process was
decreased to 38 C. Subsequently, electrodeposition coating
was performed using an electrodeposition paint, V-50,
manufactured by Nippon Paint Co., Ltd. The target amount of
deposited chemical conversion film ranged from 2 to 2.5
g/m2, and the target film thickness in electrodeposition
coating was 25 pm.
Corrosion resistance after painting was determined in
a warm salt water immersion test (SDT). A crosscut was
formed
CA 02693489 2010-01-15
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with a cutter in a sample subjected to chemical conversion
treatment and electrodeposition coating. The sample was
immersed in warm salt water (5% NaC1 at 55 C) for 10 days,
was then washed with water, and was dried. Tape peeling on
the crosscut was performed to measure the maximum peel width
on the left and right sides of the crosscut. A one-side
maximum peel width of 3.0 mm or less was considered as high
corrosion resistance after painting.
Table 2 shows the results, together with manufacturing
conditions.
Table 2
One-
Stretch
Slab Finish- Elongation
Precipitated Precipitated Amount of side
Type Coiling flangeability
heating rolling TS after Ti content V content V in
solid maximum
No of temperature after
Phase Note
temperature temperatureIMPa) prestraining for <20 nm for <20 run
solution peel
o c) (%)
steel = (*C) working:
Cc) ( (mass ppm)
(mass ppm) (mass ppm) width
Alp(%)
(nun)
1 A 1250 920 630 812 20 87 752 818 ,
340 1.7 Ferrite: 100% Example
2 B 1300 926 632 952 18 79 1580 765 231
2.5 Ferrite: 100% Example
3 C 1270 911 650 966 17 81 703 1700
380 2.2 Ferrite: 100% Example
4 C 1270 900 580 865 17 95 635 657 1350
1.2 Ferrite: 99%, Remainder: Cementite 1% Example
D 1270 917 603 1190 16 61 1700 1682
213 2.2 Ferrite: 98%, Remainder: Bainite 2% Example
6 E 1250 921 611 940 18 92 808 658 476
1.2 Ferrite: 100% Example
7 F 1250 900 590 834 20 98 727 735 540
1.4 Ferrite: 100% Example
8 G 1250 918 670 815 19 82 230 450 420
1.4 Ferrite: 100% Example
9 H 1250 920 580 802 18 93 352 167 272
1.2 Ferrite: 100% Example
I 1160 905 625 785 22 97 532 372 306
1.2 Ferrite: 100% Example
0
11 J 1250 920 630 936 18 83 863 1129
274 2.0 Ferrite: 100% Example
12 A 1250 920 480 760 19 63 150 121 934
2.0 Ferrite: 100% Comparative Example 0
IV
13 G 1250 920 720 765 18 90 220 98 330
5.2 Ferrite: 100% Comparative Example cs
14 G 1250 915 750 760 15 78 140 80 908
5.1 Ferrite: 100% Comparative Example I lf)
Lk)
K 1250 923 590 851 20 45 821 702 , 568 0.8
Ferrite: 90%, Remainder: Pearlite 10% Comparative Example
CU
co
16 L 1250 918 585 659 25 60 50 45 568
1.1 Ferrite: 100% Comparative Example
17 M 1250 918 595 850 17 40 480 353 330
0.8 Ferrite: 92%, Remainder:
Cementite 8% Comparative Example Iv
18 N 1250 920 575 765 18 75 560 540 247
1.0 Ferrite: 100% Comparative Example I (z)
H
19 0 1250 916 565 851 14 43 560 432 350
1.1 Ferrite: 100% Comparative Example i (z)
(z)
20., P 1160 921 575 653 23 75 180 324 832
1.2 Ferrite: 100% Comparative Example H
i
21 Q 1160 922 650 765 16 73 490 14 223
1.2 Ferrite: 100% Comparative Example H
22 Q 1160 920 510 782 16 50 502 220 90
1.1 Ferrite: 100% Comparative Example tn
23 R 1250 910 605 1280 13 93 2065 602
580 5.5 Ferrite: 100% Comparative Example
24 S 1250 900 610 1290 14 91 971 1890
530 5.3 Ferrite: 100% Comparative Example
29 A 1250 935 600 825 19 70 800 825 340
2.0 Ferrite: 100% Example
30 A 1260 980 580 820 19 68 802 830 355
2.1 Ferrite: 100% Example
_
31 A 1260 1020 630 826 18 73 601 824 349
2.1 Ferrite: 100% Example
32 J 1260 940 620 982 17 63 923 1120 270
2.6 Ferrite: 100% Example .
33 C 1260 960 600 983 17 65 812 1702 375
2.5 Ferrite: 100% Example
34 X 1300 965 600 1005 16 62 1205 1108
305 2.8 Ferrite: 100% Example
CA 02693489 2010-01-15
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Table 2 shows that the working examples had a TS of 780
MPa or more, Xn of 60% or more, and an SDT one-side maximum
peel width of 3.0 mm or less, indicating that the hot-rolled
steel sheets had high stretch flangeability after working
and corrosion resistance after painting.
In contrast, the comparative examples had a low TS
(strength), small Xn (stretch flangeability after working),
and/or a large SDT one-side maximum peel width (corrosion
resistance after painting).
EXAMPLE 2
Steel having a composition shown in Table 3 was melted
in a converter and was formed into a steel slab by
continuous casting. The steel slab was subjected to heating,
hot rolling, and coiling under conditions shown in Table 4
to form a hot-rolled steel sheet having a thickness of 2.0
mm.
Table 3
Composition (mass%)
Type
of
Note
steel
Si Mn P S Al Ti V Cr
W Zr
= 0.040 0.01 1.40 0.01 0.0014
0.03 0.100 0.115 0.10 ¨ Conforming steel
= 0.040 0.02 1.43 0.01 0.0015
0.03 0.104 0.105 ¨ 0.150 ¨ Conforming steel
/ 0.041 0.01 1.42 0.01 0.0014
0.03 0.102 0.105 ¨ 0.0030 Conforming steel
= 0.040 0.02 1.40 0.01 0.0014
0.03 0.101 0.115 0.20 0.140 0.0050 Conforming steel
0
0
0
0
CA 02693489 2010-01-15
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In the same way as in Example 1, the microstructure of
the hot-rolled steel sheet thus formed was analyzed to
determine the Ti content and the V content of precipitates
having a size of less than 20 nm and the amount of V in
solid solution. In the same way as in Example 1, the
tensile strength TS, the stretch flangeability after working
Xio, and the corrosion resistance after painting (SDT one-
side maximum peel width) were measured.
Table 4 shows the results.
Table 4
.
One-side
Slab Finish- Elongation Stretch Precipitated
Precipitated Amount of V
Type Coiling
maximum
heating rolling TS after
flangeability Ti content V content in solid
No of temperature
peel Phase Note
temperature temperature (MPa)
prestraining after working: for <20 nm for <20 nm solution
steel ( C)
width
( C) ( C) (%) Aio(%) (mass
ppm) (mass ppm) (mass ppm)
(mm)
25 T 1250 921 625 832 17 99 750
815 250 2.5 Ferrite: 100% Example .
26 U 1250 918 620 830 18 90 753
760 252 2.2 Ferrite: 100% Example
27 V 1250 920 621 829 17 93 753
770 250 2.0 Ferrite: 100% Example
28 W 1250 921 620 842 18 98 760
823 251 2.6 Ferrite: 100% Example
35 T 1250 940 600 835 18 , 92 780
820 240 2.2 Ferrite: 100% Example
36 T 1270 960 630 840 17 93 782
823 244 2.1 Ferrite: 100% Example
37 T 1300 980 620 837 18 95 788
830 245 2.3 Ferrite: 100% Example
n
0
1.)
m
ko
w
I
a,
m
ko
L.)
0
H
I
0
O
H
I
H
ITI
CA 02693489 2010-01-15
- 39 -
Table 4 shows that the working examples had a TS of 780
MPa or more, Xn of 60% or more, and an SDT one-side maximum
peel width of 3.0 mm or less, indicating that the hot-rolled
steel sheets had high stretch flangeability after working
and corrosion resistance after painting.
As compared with the steel sheet No. 1 (Table 2), the
steel sheets Nos. 25 to 28 and 35 to 37, which further
contained Cr, W, or Zr, had an improved TS.
Industrial Applicability
A steel sheet according to the present invention had
high strength, high stretch flangeability after working, and
high corrosion resistance after painting, and is therefore
most suitable for, for example, automobile and truck frames,
and components that require elongation and stretch
flangeability.
