DESCRIPTION
Title of Invention
High Strength, Hot Dipped Galvanized Steel Sheet
Excellent in Shapeability and Method of Production of
Same
Technical Fleld
[ ~ O O I ] The present invention relates to hlgh strength,
hot dipped galvanized steel sheet excellent ln
shapeability whicli 1s mainly suitable for auto parts and
a method of prodtlctlon of the same.
Background Art
C O O Q ~ ] Reduction of weight of thecross members, side
members, and other members of automobiles has been
considered so as to deal with the trend toward reduction
of weight for improvement of fuel economy in recent
years. In terms of materials, from the viewpoint of
securing strength and impact safety even if made thinner,
steel sheet is being made higher in strength. However,
the shapeability of materials deteriorates along with the
rise of strength, so to achieve lighter weight of the
members, steel sheet which satisfies both shapeability
and high strength has to be produced.
TOBO3J As steel sheet which achieves both shapeability
and high strength, PLT 1 discloses the art of utilizing
residual austenite and using transformation-induced
plasticity to improve the ductility, so-called residual
austen,ite steel. However, to enable residual austenite to
remain, the cooling rate after dual-phase annealing has
to be increased to prevent ferrite transformation and
pearlite'transformation and Si and A1 have to be added to
suppress the precipitation of cementite. To make the
cooling rate greater, a continuous annealing line with a
high cooling rate becomes necessary. Addi-tion of a high
content of Si impairs the plateability, while a high
content of A1 often impairs the castabiiity.
p O Q O 4 ~ PLT 2 and PLT 3 disclose so-called "dual phase
steel" which has a composite structure of low temperature
transformed phases which contain ferri-te and martensite
(hereinafter referred to as "DP steel"). This is being
widely used. DP steel exhibits a sufficient s-trengthductility
balance, through not reaching that of residual
austenite steel, so is used for relatively complicatedly
shaped chassis parts. Further, the strength of DP steel
is increasing along wi-th the trend toward lighter weight
of chasses in recent years.
g0005g For example, PLT 4 and PLT 5 disclose the art
of adding the carbide-forming elements of Nb, Tj~, and
other elements to suppress recrystallization during
annealing and utilize precipitation strengthening so as
secure a tensile strength of 780 MPa or more.
[0006] Further, PLT 6 and PLT 7 relate to composite
structure steels, where the stretch flange formability is
generally low in level, and show the art of controlling
the difference in hardnesses of the base phase ferrite
and the low temperature transformed phases so as to
improve the stretch flange formability. In these
inventions, the hardness is measvred by the Vicker's
hardness.
[0007] However, in the case of DP steel with :I tensile
strength of '780 MPa or more, the grain size of the
microstructure is small and the hardness cannot be
evaluated by the Vicker's hardness. Therefore, art which
evaluates properties by the nano hardness which is
measured using the newly developed art of nano
indentation is disclosed in PLT 8. With this art, the
ratio of hardnesses of the ferrite and the low
temperature transformed phases is defined in accordance
with the ferrite fraction. Due to this, the bending
properties are improved.
Citations List
Patent Literature
[ O O O ~ ] PLT 1: Japanese Patent Publication (A) No. 6-
145788
PLT 2: Japanese Patent Publication (A) No. 10-147838
PLT 3: Japanese Patent Publication (A) No. 2002-
363695
PLT 4: Japanese Patent Publication (A) No. 2009-
144225
PLT 5: Japanese Patent Publication (A) No. 2002-
363685
PLT 6: Japanese Patent Publication (A) No. 2009-
191360
PLT 7: Japanzse Patent Publication (A) No. 2009-
167475
PLT 8: Japanese Patent Publication (A) No. 2009-
167467
PLT 9: Japanese Patent Publication (A) No. 2010-
65316
Summary of Invention
Technical Problem
[BOB91 When producing DP steel utilizing microalloy
elements, the precipitation behavior of Ti, Nb, and other
microalloy carbides has an effect on the material
quality. That is, the material quality sometimes
fluctuates diie to the production conditions of vieel
sheet, in particular the annealing conditions. The
fluctuations in the material quality in this case mainly
appear in the yield strength and the stretch flange
formability.
[OBlOg When the yield strength fluc-tuates, there are
the problems that the springback behavior changes after
press forming and the dimensional precision of chassis
parts falls. For this reason, defective parts occur and
correction off-line becomes necessary.
If the. stretch flange formability fluctuates, there is
the problem of cracking at the sheared parts which are
produced by blanking at the time of press forming.
I O Q l l ] I n the past, tne element which has been used
most as a microalloy is Nb. Nb, even in the solid
solution state, retards ferrite grain growth or
recrystallization due to the solute dragging effect and
contributes to strengthening due to non-recrystallized
ferrite and increased grain fineness. Furthermore, when
precipitated as carbides, there is an effect of
improvement of strength due to precipitation
strengthening. Due to these reasons, Nb has been used for
improving strength.
Further, if adding B to this, the solute dragging effect
of Nb is improve$ and the strength raising effect becomes
larger.
[001231 However, with addition of Nb and further with
addition of Nb and B, the effect of retarding
recrystallization and the effect of suppressing grain
growth are large, so a high annealing temperature becomes
necessary. For this reason, in the general temperature
range of continuous annealing, that is, 720°C to 800°C in
range, recrystallization is not completed and the
dependency of the tensile strength or other aspects of
material quality on the annea1ing.temperature ends up
become larger.
[00%.3] Fur-ther, the art which is disclosed in PLT 8
for restriction of the ratio of the nano hardnesses of
the ferrite phase and low temperature transformed phases
to improve the bendability just defines the ratio of the
acerage hardnesses. For this reason, even if there is a
phase of a high hardness in the low temperature
transf,ormed phases, sometimes this ends up being included
in the a-verage value with the surroundings. When there is
such a high hardness low temperature transformed phase,
this becomes a cause of fluctuations of stretch flange
formability and further becomes a cause of fluctuations
of the tensile characteristics, so this is a problem.
[0014] PLT 9 discloses a composite structure steel
sheet which has ferri-te as a main phase and bainite and
martensite low temperature transformed phases as
secondary phases wherein the balances of TS-EL and TS-h
(measure for evaluation of stretch flange formability)
are good. This discloses the deliberate addition of Ti
and Nb, con-trol of the ratios of composition of the
secondary phases, and control of the hardness of the base
material structure. However, in this case, an Ac3
temperature or higher annealing temperature becomes
necessary, so the annealing temperature dependency is
large.
[QQ%51 In this way, keeping the material quality from
fluctuating due t,o the manufac-turing conditions i.s an
important requirement for DP steel. In particular, DP
steel which does not fluctuate in material quality under
high productivity annealing conditions, for example, even
in the general temperature range in continuous annealing,
that is, 720°C to 8OO0C in range, or at most annealing at
the Ac3 temperature or less, is being sought.
Solution to Problem
[OSISl The inventors engaged in in-depth studies to
solve the above problems and as a,result discovered that
by adding Ti, which is smaller in the effect of retardjng
recrystalliz,ation and grain growth compared with Nb and
which enables recrystallization in the general
temperature range in continuous annealing, that is, 720°C
t~ 8OO0C in range, and by limiting the amounts of addition
of Nb and B, it is possible to suppress fluctuations in
the material quality.
~ 0 0 1 7 ~ ' That is, they discovered that even if annealing
at a temperature range of 720°C to a temperature of the
iower ofr85O0C or the Ac3 temperature (below, referred to
as the "easy annealing temperature region") without
heating to the annealing Ac3 temperature, it is possible
to suppress fluctuations in material quali-ty while
obtaining predetermined properties.
COOIS] Further, the inventors discovered that by
optimizing the heating rate and the cooling pattern at
the time of annealing, it is possible to con-trol the
ferrite phase rate, the grain size of the low temperature
transformed phases, the ratio of the average values of
nano hardnesses of the ferrite phase and the low
temperature transformed phases, and fluctuations in
hardnesses of the low temperature transformed phases.
[0019] Due to these findings, the inventors discovered
that it is possible to produce high strength, hot dipped
galvanized steel sheet with little fluctuations in
material quality and completed the present invention. The
gist of the present invention is as follows:
600201 [I] High strength, hot dipped galvanized steel
sheet characterized by containing, as ingredients of the
steel, by mass%, C: 0.05 to 0.1%, Si: 0.1 to 1.0%, Mn:
2.0% to 2.5%, Al: 0.02 to 0.18, Ti: 0.01 to 0.05%, Cr:
0.1 to 1.0%, Sn: 0.0010 to 0.1%, and a balance of Fe and
unavoidable impurities, having a microstructure comprised
of low temperature transformed phases of a ferrite phase
fraction of 70 to 90% and a balance of martensite, having
an average grain size of the low temperature transformed
phases of 0.1 to 1 pm, having a ratio of average nano
hardnesses of the ferrite phase and the low temperature
transformed phase of 1.5 to 3.0, and having a ri~ino
hardness of the low temperature transformed phases at 80%
or more of the measurement points of 1 to 5 times the
average nano hardness of the ferrite phase.
100211 [2] A method of production of high strength,
hot dipped galvanized steel sheet characterized by
heating a slab which has the steel ingredients as set
forth in [I] to 1000 to 1350°C, then hot rolling at a
final rolling temperature Ar3 or more, coiling at 600°C or
less, pickling, cold rolling ar a rolling rate of 30 to
70%, and, after that, heat treating while making a
temperature of '720°C to a temperature of the lower of
850°C or the Ac3 temperature the annealing temperature,
during which heating in the temperature range from a.t
least 600°C to the annealing temperature by a 0.5OC/sec to
G°C/sec heating rate, holding at the annealing temperature
for 10 sec or more, then cooling in at least the
temperature range of the annealing -temperature to 650°C by
a cooling rate of 5'C/sec or more, fur-ther cooling in at
least the temperature range of 600°C to 500°C by a cooling
rate of 3'C/sec or less, then performing hot dip
galvanization or hot dip galvannealization. The annealing
temperature is made the 850°C or Ac3 temperature or less
because if heating over these temperatures, the steel
sheet strength rapidly falls and the runnability at the
annealing step becomes poor.
Advantageous Effects of Invention
100221 According to the present invention, in the
continuous annealing step, even if annealing in the
temperature range of the gencral annealing temperature of
720°C to a temperature of the lower of 850°C or the Ac3
temperature (easy annealing temperature region), it is
possible to provide high strength, hot dipped galvanized
steel sheet which has little fluctuation in material
quality, which is excellent in shapeability, arid which
has predetermined properties. In particular, a remarkable
effect is exhibited in 780 MPa or more high strength, hot
dTpped galvanized steel sheet.
Brief pescription of Drawings
[0024% FIG. 1: FIG. 1 is a graph which shows the
relationship between the recrysta.llization rate and
annealing temperature of Ti-containing steel, Nbcontaining
steel, and Nb-B-containing steel.
FIG. 2: FIG. 2 is a micrograph which shows the
structure of steel sheet used when finding the area rate
of the low temperature transformed phases in the
invention examples.
Description of Embodiments
[88%4] The basic thinking of the present invention
will be explained.
The steel sheet according to the present invention limits
the addition of carbide-forming elements to Ti, does not
add the conventionally often used Nb, and, further,
limits the amount of addition of B - which has a great
effect on recrystallization.
[0025] Further, the method of production of steel
sheet according to the present invention is charac-terized
by optimizing the heating rate at the time of annealing
and the cooling rate after annealing so as to control the
ferrite phase rate, the grain size of the low temperature
%ransformed phases, the ratio of the average values of
the nano hardnesses of the ferrite phase and the low
temperature transformed phases, and the fluctuations in
the hardnesses of the low temperature transformed. phases.
Due to this, they discovered that high strength hot
dipped galvariized steel sheet with little fluctuations in
material quality can be produced.
10026j First, the reasons for limiting the carbideforming
elements which contribute to the
recrystallization of ferrite and precipitation
strengthening to Ti and for limiting the contents of the
carbide-forming elements will be shown below.
[+I0274 Nb is an element effective for suppression of
recrystallization and precipitation strengthening. The
precipitation behavior at the time of hot rolling greatly
depends on the coiling temperature. Further, even in the
case of fine precipitation at the time of heating of the
annealing step, the effect of retardation of
recrystallization is large, so the material quality of
the annealed steel sheet is believed to greatly depend on
the annealing temperature.
[8828] The same is true for the V, W, Mo, and Zr which
are used as other carbide-forming elements. Further, the
dependency of the carbide precipitation on the coiling
temperature in hot rolling and the dependency on the
heating rate and dependency on the annealing temperature
in the annealing process differ, so become causes of
fluctuation of the material quality.
[00%9] Ti, compared with Nb, has a small effect of
retarda-tion of recrystallization or grain growtl~ due to
the solute dragging effect or the effect of precipitation
strengthening due to carbides. For this reason, in the
temperature range where manufacture is easy in general
continuous annealing, that is, 720°C to 800°C in range,
the dependency of the tensile strength and othe~ facets
of material quality on the annealing temperature becomes
small. A conceptual view is shown in FIG. 1.
[0030] As shown in FIG. 1, in Ti-containing steel, it
is believed that recrystallization starts at a
temperature in the ferrite region and that the austenite
transformed to when reaching the dual phase region of
ferrite and austenite retards the recrystallization. If
raising the temperature after that, the recrystallization
ends. Due to this, with Ti-containing steel, compared
with Nb-containing steel and Nb-B-containing steel, it is
possible to obtain a comparatively low temperature,
stable recrystallized state. That is, it is be:liLeved
possible to reduce the fluctuations in material quality.
Due to the above reasons, the carbide-forming elements
aye limited to only Ti.
[0031g Further, B retards the ferrite transformation
and pearlite transformation at the time of cooling after
annealing. Therefore, this is an effective element for
obtaining a composite structure. However, the effect of
suppression of transformation is large, so the steel
sheet after hot rolling is hard and cold rolling
sometines becomes difficult. Further, as shown by the
example of the Nb-B-containing steel of FIG. 1,
recrystallization in the easy annealing temperature
region is suppressed, so the annealing temperature lias to
be raised. For this reason, the amount of addition of B
was limited. A conceptual view of the retardation of
recrystallization in Nb-B-containing steel is shown
together in FIG. 1.
[0032] To control the fine precipitation of Ti during
the heating of the annealing step, limitation of the
amount of addition of Ti and the hot rolling conditions,
in particular the coiling temperature, is effective.
Furthermore, by limiting the heating rate by the heating
during the annealing step, the retardation of the
recrystallization is suppressed and -the fluctua-tions in
the material quality can be reduced.
COO331 Ti precipitates as TiN in an amount
corresponding to mainly the amount of addition of N at
the time of heating for hot rolling. The remaining Ti
precipitates as Tic at the time of coilinq, so the
coiling temperature is limited to suppress fine
precipitation. The Ti which did not form any precipitates
at the time of hot rolling, that is, the solid. solution
Ti, is believed to finely precipitate as Tic at the time
of 'heating in the annealing step or remain in the form of
solid solution Ti so as to suppress recrystallization and
grain growth.
[O034] Next, control of the low temperature
transformed phases will be explained. The grain size of
the low temperature transformed phases, the ratio of
hardnesses of the low temperature transformed phases and
ferrite phase, and the ranges of their fluctuations were
controlled. In DP steel, the tensile strength is greatly
affected by the strength of the low temperature
transformed phases. That is, when the hardnesses of the
low temperature transformed phases are high, the tensile
strength becomes high. For this reason, fluctuations in
the hardnesses of the low temperature transformed phases
become the cause of fluctuations in the tensile strength.
The hardnesses of the low temperature transformed phases
depend on the concentration of carbon in the austenite at
the time of annealing. Further, if the amount of carbon
fluc-tuates, the transformation expansioii rate fluctuates
and the amount of movable dislocations which are
introduced into the nearby ferrite is affected. For this
reason, by limiting the ratio of hardnesses of the low
temperature transformed phases and the ferrite phase and
the ranges of their fluctuation, it is possible to
suppress the fluctuation of the yield strength.
[0035] The ratio of hardnesses of the low temperature
transformed phases and the ferrite phase also has an
effect on the strstch flange formability. With DP steel,
voids are formed from near the low temperature
transformed phases and act as the starting points of
cracks. If the ratio of hardnesses of the low temperature
transformed phases and the ferrite phase is large, even
if the strain is small, voids easily form. From the
viewpoint of this stretch flange formability, the ratio
of hardnesses is preferably small.
&0036] However, if the ratio of hardnesses is
excessively small, the expansion of volume when the low
temperature transformed phases transform will be small,
so the amount of movable dislocations which are
introduced into the ferrite becomes smaller. In this
case, the low yield ratio characterizing DP stei'L can no
longer be realized.
/8033] The hardnesses of the low temperature
transformed phases depend on the concentration of carbon
in the austenite. If the distribution of carbon in the
austenite becomes excessively uneven, the fluctuations in
hardnesses of the low temperature transformed phases
become greater and, along with this, the fluctuations in
th,e yield strength and stretch flange formability become
greater. For this reason, control of the ex-tent of
fluctuation of the hardnesses of the low temperature
transformed phases is important in controlling
fluctuations in material quality.
[I30381 The low temperature transformed phases are
preferably fine in graln size and dispersed in large
amounts. The reasons are that voids are not locally
formed at the time of stretch flange formation, so this
is advantageous, and that fine dispersion results in
uniform introduction of movable dislocations into the
ferrite.
EOO391 The ratio of hardnesses of the low temperature
transformed phases and the ranges of fluctuation and
grain size can be controlled by the heating rate and the
cooling rate in the annealing step. The thinking is shown
below.
[004O] First, the heating rate will be explained. In
the heating step of annealing, melting of the iron
carbides, recovery of the ferrite, and recrystallization
occur near 600°C or more and the ferrite transforms to
austenite at the Acl transformation point near 700°C or
more. The melting of iron carbides is promoted by
lowering the heating rate, while the distribution of
carbon is made uniform. If ferrite is transformed to
austenite, the recrystallization of ferrite is
suppressed. For this reason, by limi-ting the heating rate
in the temperature region from 600°C to the annealing
temperature, it is possible to control the
recrystallization rate.
[0041] The fractions of ferrite and austenite are
determined by the annealing temperature, - while carbon
etc. concentrates at the austenite. Further, due to
limitations on the heating rate, the amount of addition
of Ti,' and the coiling temperature at hot rolling, the
recrystallization of ferrite is con%rolled and the ratio
of hardnesses of the ferrite and low temperature
transformed phases is held at a suitable range.
[0042] Next, with cooling after annealing, the sheet
is cooled relatively fast in the temperature range from
the annealing temperature down to 650°C so as to increase
the nucleation sites of transformation and make the low
temperature transformed phases finer. Further, by cooling
relatively slowly over the temperature range from 600°C to
500°C, it is possible to reduce the fluctuations of the
amount of carbon in the austeni-te which is distributed by
ferrite transformation.
EQQ433 Next, the reasons for limitation of the
specific conditions will be explained.
First, the limitations of the chemical ingredien-ts will
be explained. Nots that unless particularly indicated,
r, g ,, means mass%.
10044J C: C is an element which can raise the strength
of the steel sheet. However, if less than 0.05%, the
hardness of the mainly martensite low temperature
transformed phases becomes lower, so securing the 780 MPa
or more tensile strength becomes difficult. On the other
hand, if over 0.1%, securing the spot weldability becomes
difficult. For this reason, the range is limited to 0.05
to 0.1%. To reliably obtain this effect, the lower limit
value is preferably made 0.06%, more preferably 0.07%, if
pos,sible preferably 0.075%. Further, the upper limit
value is preferably made 0.095%, if possible preferably
0.09%.
[QQ45] Si: Si is a strengthening element and is
effective for raising the strength of the steel sheet.
However, if less than 0.1%, the drop in shapeability due
to deterioration of the elongation - becomes remarkable.
Further, if over I%, the plating wettability falls.
Therefore, the Si content is restricted to 0.1 to 1.0% in
range.. To obtain this effect reliably, the lower limit
value is preferably made 0.25%, more preferably 0.3%, if
possible 0.4%. Further, the upper limit value is
preferably 0.8%, if possible 0.6%, more preferably 0.5%.
For a continuous hot dip galvanization line which has an
all radiant tube type heating furnace, 0.4 to 0.5% is
most suitable.
E0046J Mn: Mn is a strengthening element and is
effective for raising the strength of the steel sheet.
However, if less than 2.0%, obtaining a 780 MPa or higher
tensile strength is difficult. If conversely large, it
aids co-segregation with P and S and invites a remarkable
deterioration in the bendability, elongation, and hole
expandability, so 2.5% is made the upper limit. To
reliably obtain this effect, the lower limit value is
preferably 2.1%, more preferably 2.2%. The v.pper iimit
value is preferably 2.4%, more preferably 2.3%.
%0047g Ti: Ti is an important element which
contributes to the rise in strength of steel sheet by
suppression of growth of ferrite crystal grains and
therefore grain size reduction strengthening and
dislocation strengthening. It hardens the main phase
ferrite and lowers the difference in hardnesses of the
strengthening phases, that is, the mainly martensite low
temperature transformed phases and ferrite phase so as to
improve the bendability and hole expandability. These
effects cannot be obtained if less than 0.01%, so the
lower limit value was made 0.01%.
On the other hand, if containing over 0.05%, the
precipitation of carbonitrides becomes greater and the
shapeability deteriorates, so the upper limit value is
made 0.05%. To make the effect reliable, the lower limit
value is preierably made 0.015% and more prefer.ibly is
made 0.02%. The upper limit value is preferably 0.04%,
more preferably 0.03%. If making the tensile strength 780
MPa or more and reducing the fluctuations in the yield.
strength, the lower limit value should be limited to
0.02% and the upper limit value to 0.03%.
[0048] Cr: Cr is a strengthening element and is
important for improvement of hardenability. It is an
austeni-te former, so is an element essential for securing
the austenite fraction at a low temperature. If less than
0.1%, these effects cannot be obtained, so the lower
limit value was made 0.1%. Conversely, if containing over
I%, the strength excessively increases, so the upper
limit value was made 1%. Preferably, the content is made
0.2 to 0.8%, more preferably 0.3 to 0~7%.
[0049] Al: A1 promotes formation of ferrite and
improves the ductility, so may be added. Further, it may
also be utilized as a deoxidizing material. The effect is
not exhibited if less than 0.02%, so tb.e lower limit was
made 0.02%. However, excessive addition forms Al-based
coarse inclusions and becomes a cause of surface damage
and deterioration of hole expandability. Due to this, the
upper limit of addition of A1 was made 0.1%. Preferably,
the content is 0.04 to 0.09%, more preferably 0.05 to
0.08%.
60050l P: P tends to segregate at the center part of
thickness of the steel sheet and causes the weld zone to
become brittle. For this reason, the smaller the amount
the better, but zero is rather better. If over 0.03%, the
embrittlement of the weld zone becomes remarkable, so the
suitable range is limited to 0.03% or less. The lower
limit value of P is not particularly set, but making the
content less than 0.0001 mass% would be disadvantageous
economically, so this value is preferably made the lower
limit value. That is, the content.which is allowed as an
unavoidable impurity is made 0.03% cr less.
[0051] S: S has a detrimental effect on the
weldability and on the manufacturability at the time of
casting and time of hot rolling. For this reason, the
smaller the amount the better, but zero is rather better.
be to this, the upper limit value was made 0.01 mass% or
less. The lower limit value of S is not particularly set,
but making the content less than 0.0001% is
disadvantageous economically, so this value is preferably
made the lower limit value. That is, the content which is
allowed as an unavoidable impurity is made 0.01% or less.
[:0052] N: N forms coarse nitrides and causes
deterioration of the bendability and hole expandability,
so the amount of addition must be kept down. Zero is
rather better. This is because if N is over 0.01%. this
tendency becomes remarkable, so the rang-e of N conterlt is
made 0.01% or less. In addition, this becomes a cause of
occurrence of blowhoies at the time of welding, so the
less the better. The lower limit is not particularly set.
The effect of the present invention is exhibited, but
making the N content less than 0.0005% invites a major
increase in the manufacturing costs, so this is the
substantive lower limit. That is, the content which is
allowed as an unavoidable impurity is 0.01% or less.
[0053] Nb: Nb ;s an element which has an effect of
suppression of recrystallization and is effective for
refining the ferrite and strengthening the ferrite phase
due to precipitation strengthening. However, it
precipitates as NbC during rolling and in the coiling
step at the time of hot rolling and during the heating of
the annealing step to thereby affect the precipitation
strengthening and suppression of recrystallization. Even
as solid solution Nb, due to the solute dragging effect,
it affects the suppression of recrystallization, so
greatly affects the strength. Therefore, it is
susceptible to the effects of the production process and
becomes a cause of fluctuations in'material quality.
Addition is therefore not preferred. Therefore, in the
present invention, this is not deliberately added. Even
if present, the content is preferably limited to 0.0010%
or less. This limitation is set because even in the case
of basically zero addition, considering -- the case that
utilization of scraps results in this element ending up
being included, it is preferable to control the content
to this limitation or less. That is, the content which is
allowed as an unavoidable impurity is made 0.0010% or
less.
[0054] V, W, Mo, and Zr: These carbide-forming
elements are characterized by greater difficulty in
forming precipitates compared with Ti and Nb. If these
elements are added, they become causes of fluctuation of
material quality since the precipitation behaviors differ
for the respective elements and the dependency of the
carbide precipitation on the coiling -temperature and the
dependency on the heating rate or dependency on the
annealing temperature change in the annealing step. For
this reason, addition is not preferable. Therefore, in
the present invention, V, W, Mo, and Zr are not
deliberately added. Even if present, their contents are
preferably restricted to 0.0010% or less. This ;imitation
is set because even in the case of basically zero
addition, considering the case that utilization of scraps
results in these elements ending up being included, it is
preferable to corirol the contents to this limitation or
less. That is, the contents which are allowed as
unavoidable impurities are made 0.0010% or less.
100551 B: B is an element which makes the
quenchability increase and is also effective for
suppression of recrystallization. However, due to the
addition of B, the strength of the hot rolled steel sheet
rises and the cold rollability falls. Further, to
suppress the recrystallization of ferrite, the annealing
temperature has to be raised. Zero addition is
preferable. Therefore, in the present invention, this is
not deliberately added. Even if present, the content is
preferably limited to 0.0001% or less. This limitation is
set because even in the case of basically zero a~ddition,
considering the case that utilization of scraps results
in this element ending up being included, it is
preferable to con-trol the content to the limitation or
less. That is, the content which is allowed as an
unavoiclable impurity is made 0.0001% or less.
60056] Sn: Sn improves the plating adhesion at the
time of hot dip galvanization and further has the effect
of promoting alloying. The effect is not exhibited if
less than 0.0010%, so the lower limit was made 0.0010%.
Further, if excessively added, the hot workability of the
slab falls, so the upper limit was made 0.1% or less. To
reliably obtain this effect, the lower limi-t value is
preferably made 0.002% and the upper limit value is
preferably made 0.03%. Furthermore, the lower 1imi.t value
is more preferably made 0.005% and the upper limit value
is more preferably made 0.01%.
[00.57% As other elements, Ca or an REM may also be
added for control of the form of the sulfides. Further,
sometimes Ni, Cu, and other elements are contained as
unavoidable impurities, but any content is posslbie so
long as the content does not have any effect of the
properties of the present invention. The contents of
these elements are preferably, as a general measure,
0.05% or less for the respective elements.
[0058] Next, the reasons for limitation of tile
microstructure will be explained.
The ferrite phase fraction is made 70 to 90% and the
balance is made martensite and other low temperaturc
transformed phases. By making this ratio, a tensile
strength of 780 MPa or more and a predetermined ductility
are secured. If the ferrite phase fraction is less than
70%, the ductility due to ferrite cannot be secured. If
the ferrite phase fraction is over 90%, the content of
the low temperature transformed phases is small, so the
tensile strength falls below 780 MPa. The ferrite phase
fraction is preferably 75 to 88%, more preferably 80 to
85%.
600591 The low temperature transformed phases are made
to contain martensite so as to enable use of martensite
transformation to cause movable dislocations to be
introduced into the ferrite phase and the yield point to
fall and to secure a yield ratio of 0.7 or less.
10060] The low temperature transformed phases are
preferably finely dispersed in large amounts. Due to
this, not only does the stretch flange formability become
excellent, but also the introduction of movable
dislocations in the ferrite phase becomes uniform.
However, if the average grain size of the low temperature
transformed phases is less than 0.1 pmcmt,h e amount of
introduction of movable dislocations into the ferrite
becomes small and the yield ratio exceeds 0.7. For this
reason, the lower limit of the average grain size of the
low temperature transformed phases was made 0.1 pm.
Further, if the average grain size of the low temperature
transformed phases is excessive, the stretch flange
formability deteriorates, so the upper limit was made 1
pm. To reliably obtain this effect, the average grain
size of the low temperature transformed phases is more
preferably 0.4 to. 0.8 pm in range. More preferably, it
may be made 0.5 to 0.7 pm.
Note that, for the method of measurement of the grain
size of the ferrite phase fraction or low temperature
transformed phases, this can be measured based on the
LePera method which is described in the [Average Grain
Size of Low Temperature Transformed Phases] in the later
explained invention exampies.
[0061] The ratio of the average nano hardnesses of the
ferrite phase and the low temperature transformed phases
may be made 1.5 to 3.0 (defined as the average nano
hardness of the low temperature transformed
phases/average nano hardness of the ferrite phase. The
nano hardness is measured at a position of a depth of
about 1/4 of the thickness from the steel sheet surface).
If the ratio of hardnesses exceeds 3.0, the stretch
flange formability deteriorates. Further, if the ratio of
hardnesses is less than 1.5, the concentration of carbon
at the low temperature transformed phases becomes
insufficient and the introduction of movable dislocations
in-to the ferrite due to the volume expansion of the
martensite transformation becomes insufficient. For this
reason, the low yield ratio characterizing DP steel can
no longer be secured. The lower limit of the ratio of
average nano hardnesses is more preferably 1.7, still
more preferably 1.9. Further, the upper limit of the
ratio of the average nano hardnesses is more preferably
2.8 and still more preferably 2.5.
[0062] The "nano hardness" is the ultrasmall loading
hardness using a pyramidal indenter which is defined in
JIS Z 2255. The measurement load was made 1 mN. The nano
hardness sometimes fluctuates due to the measurement
load. In the case of the present invention steel, the
measurement load is optimally 1 mN in view of the
relationship of the grain size of the low temperature
transformed phases and the indentation. The nano hardness
is defined by the value measured at this load. The
average nano hardness is found from the results of
measurement of a minimum of 30 points or more, preferably
100 points or so.
[0063] Regarding conventional composite structure
steel, the above-mentioned PLT 6 and PLT 7 disclose the
results of the ratio of hardnesses based on the Vicker's
hardness. However, they do not disclose anything, like in
the present invention, regarding the effects of the ratio
of nano hardnesses between microstructures on the stretch
flange formability. The Vicker's hardness is measured by
the size of the indentation after removal of load from
the indenters, but with the nano hardness, the hardness
is found by the depth of penetratiorl of the indenter in
the load state. For this reason, this features no
deformation seen due to the elastic recovery oc,,urring in
measurement of the Vicker's hardness. That is, the nano
hardness and the Vicker's hardness clearly differ in
msasurement methods. For this reason, the effect of the
ratio of the nano hardnesses of the ferrite phases and
low temperature transformed phases on the stretch flange
formability can be said to first appear in composite
structure steel of a fine structure.
&0064] , Here, it was learned that when 80% or more of
the nano hardness measurement points of the low
temperature transformed phases have nano hardnesses
within the range of 1 to 5 times the average nano
hardness of the ferrite phase, the stretch flange
formability does not deteriorate. In other words, this is
because if 20% or more of the nano hardness measurement
points of the low temperature transformed phases have
nano hardnesses over 5 times the average nano hardness of
the ferrite phase, the density of movable dislocations
which are introduced into the ferrite near the low
temperature transformed phases becomes higher and the
fluctuation of the yield strength becomes larger.
Therefore, 80% or more of the nano hardness measurement
points of the low temperature transformed phases have a
nano hardness of 5 times or less of the average nano
hardness of the ferrite phase. Further, when 20% or more
of the nano hardness measurement points of the low
temperature transformed phases have less than 1 time the
average nano hardness of the ferrite phase, there is
little expansion of volume in the martensite
transformation near the low temperature transformed
phases and there are less movable dislocations which are
introduced into the ferrite. In this case as well, the
fluctuations in the yield strength become greater.
Therefore, it is prescribed that 80% or more of the
measurement points of the nano hardness of the low
temperature transformed phases have a nano hardness of 1
time or more of the average nano hardness of the ferrite
phase. When the tensile strength is made 780 Mi';! or more
and the fluctuation of the yield strength is reduced, the
value may be set to 90%. Preferably, the value is made
32% or more. Note that, the nano hardness of the low
temperature transformed phases is measured at leas-t at 10
points or more, if possible 20 points or more.
[0t365:! Next, the limitation of the tensile properties
will be explained. The yield ratio was made 0.7 or less
because if the above ingredients and microstructure are
formed, the result becomes UP steel and this is a
condition which shows the low yield ratio characterizing
DP steel.
[0066] Ten slabs which were cast by the above
ingredients were processed to produce hot dipped
galvanized steel sheet under the conditions of the above
[2] including annealing in the easy annealing temperature
region. The difference of the maximum value and the
minimum value of the yield strengths of the 10 steel
sheets was defined as the fluc-tuation of the yield
strength. When using the chemical ingredients and
microstructure of the above [I], this value may be made
60 MPa or less.
E00673 Regarding the tensile strength, by making the
chemical ingredients and microstructure the ones shown in
the above [1], it is possible to obtain a 780 MPa or more
tensile strength.
~0068g The hot dip galvanization may be the usual hot
dip galvanization or may be hot dip galvannealization.
The hot dipped galvanized steel sheet which is shown in
the above [I] may be produced by any me-thod of production
so long as the features of the chemical ingredients and
the microstructure are in the ranges which are shown in
the above [l]. However, if using the method of production
which is shown in the above [2], the sheet can be easily
produced. Due to this, the method of production of the
same will be explained.
608691 First, the conditions of the hot rolling will
be explained. The slab heating temperature was iiiade 1000
to l35O0C. This is because if less than 1000°C, the
rolling load becomes higher and, due to the drop in
tzmperature before the final rolling, the prescribed
finishing temperature cannot be secured. Further, if over
1350°C, a large amount of scale is formed and causes scale
defects.
[0070] The final rolling temperature was made Ar3 or
more because when malcing the finishing temperature a
temperature lower than this, transformation occurs during
the rolling, the rolling load greatly fluctuates, and
mis-rolling is caused. Further, at the locations where
transformation occurs, the grain size becomes coarser,
t-lie microstructure after cold rolling and annealing
becomes uneven, and fluctuations in material quality are
caused.
[0071] The coiling temperature is made 600°C or less
because by malting it this temperature or less, the
carbide-forming element Ti remains as is in the solid
solution state and contributes to the grain size
reduction strengthening and dislocation strengttieiling.
Further, there are also the effects that the fluctuations
in ma-terial quality of the hot rolled sheet strength in
the longitudinal direction of the coil become smaller and
the fluctuations in the sheet thickness at the time of
cold rolling become smaller. Furthermore, at a cooling
temperature which is over 60OoC, coarse carbides are
formed and it becomes difficult for carbides to dissolve
into the austenite during the annealing, so the ratio of
the nano hardnesses falls and the low yield ratio of the
present invention steel can no longer be realized.
[00'721 The total rolling rate of the cold rolling
(hereinafter the total rolling rate of cold rolling being
referred to as just the "rolling rate") is made 30 to
70%, while the draft per pass is made 30% or less. If the
rolling rate is less than 30%, the structure after
annealing becomes coarse and the lim'tation of the grain
size of the low temperature transformed phases which is
shown in the above [I] cannot be secured, so the lower
limit is made 30%. Further, when the rolling - rate exceeds
70%, the drive force behind the recrystallization becomes
greater and recrystallization is promoted, so it becomes
difficult to secure non-recrystallized ferrite and the
strength falls. Therefore, the upper limit was made 70%.
[00'731 Further, when the draft per pass exceeds 30%,
the strong shear band can become thin and the strain near
the shear band becomes larger, so the strain profile
inside of the steel sheet becomes uneven. At the time of
annealing, the ferrite grain size in the high strain
region becomes smaller, so the uniformity of the
structure at the inside of the steel sheet falls.
[0074] Furthermore, small grain size ferrite has a
high drive force of grain growth, so the size is strongly
affected by the annealing temperature and the
fluctuations in the yield strength at. the time of
manufacture become greater.
[08753 Therefore, if making the draft per pass 30% or
less, it is possible to suppress the formation of a
strong shear band and possible to make the buildup of
strain in the steel sheet uniform. If making the draft
per pass preferably 25% or less, more preferably 20% or
less, most preferably 15% or less, it becomes possible to
make the bui~ldup of strain uniform.
600963 The annealing is preferably performed on a
continuous hot dip galvanization line. The limitations in
the temperature control at that time will be explained.
[0077] The heating rate is preferably made an average
heating rate of 0.5 to G0C/sec, preferably 0.5 to 4OC/sec,
for at least the temperature range from 600°C to the later
explained annealing temperature. The average heating rate
used is the value of the annealing temperature minus 600°C
divided by the time for reaching the annealing
temperature from 600°C (=(annealing tempera-ture-
600°C) / (time required from 600°C to annealing
temperature)). When the average heating rate is high, the
time for melting the iron carbides is insufficient and
the distribution of carbon inthe steel sheet becomes
uneven. Further, if the recrystallization of ferrite also
becomes insufficient, the strength becomes excessive and
the strength-duc-tility balance falls. Therefore, the
heating rate of the upper limit was made b°C/sec,
preferably 4'C/sec. Further, when the heating rate becomes
less than 0.5'C/secr ferrite grain growth proceeds, so the
effect of grain size reduction strengthening cannot be
expected, the strength becomes insufficient, and,
further, the required annealing line length also becomes
excessive, so the result is not economical, therefore the
lower limit was made 0.5"C/sec.
[00781 The annealing is performed by holding the sheet
at the annealing temperature of 720°C to the -temperature
of the lower of 850°C or the Ac3 temperature in
temperature range, preferably 740 to the temperature of
the lower of 800°C or the Ac3 temperature in temperature
range, for 10 seconds or more. When the annealing
temperature is less than 720°C, the amount of austenite
becomes insufficient, the tensile strength becomes less
than 780 MPa, further, the mainly martensite low
temperature transformed phases become higher in lhardness,
and the range of the above [I] is not satisfied.
Therefore, the lower limit was made 720°C. Further, by
making the lower limit of the annealing temperature 740°C,
a sufficient austenite fraction is secured and the
strength-ductility balance and stretch flange formability
become excellent.
[00793 On the other hand, excessive high temperature
heating invites a rise of costs, so not only is this not
preferable economically, but also the sheet shape at the
time of high temperature running becomes inferior, the
lifetime of the rolls is reduced, and other trouble is
caused, so upper limit of -the peak heating
temperature is made the temperature of the lower of 850°C
or the Ac3 temperature. Further, if the annealing
temperature is over 850°C, the oxides which form at the
steel sheet surface are picked up by the hearth rolls.
Sometimes pitting causing indentations at these steel
sheets occurs. Further, in the present invention, due to
the addition of Ti, recrystallization is possible even at
a temperature lower than the Ac3 temperature, so there is
no need to raise the temperature to the Ac3 or more (see
FIG. 1) .
10080g Therefore, the upper limit of the annealing
temperature is preferably made the temperature of the
lower of 850°C or the Ac3 temperature. The temperature of
the lower of 800°C or the Ac3 terr~perature is more
preferable. The heat treatment time at this temperature
region has to be 10 seconds or more so as to melt the
iron carbides. If shorter than this time, not only does
the fluctuation of the low temperature transformed phases
become greater, but also the grain size becomes
excessively fine. On the other hand, if the heat
treatment time becomes more than 600 seconds, a rise in
cost is invited, so this is not preferable economically.
[OOSl% Regarding the cooling conditions, the
temperature rangc from at least the annealing temperature
to 650°C is cooled by a cooling rate of 5OC/sec or more,
preferably 7'C/sec or more, and the temperature range from
at least 600°C to 500°C is cooled by a cooling rate of
3'C/sec or less, preferably Z°C/sec or less.
[0082% First, the cooling rate of the temperature
range from the annealing temperature to 650°C is made
faster to suppress the ferri-te transforma-tion at 650°C or
more. Due to overcooling of the ferrite, the nucleation
sites of the ferrite transformation become greater, the
ferrite becomes finer, and the grain size of the
austenite which remains at the grain boundary also is
observed to hecome finer. If making this cool in,^ rate
less than 5"C/sec, ferrite transformation occurs at a high
temperature and as a result the limitation of the average
size of the low temperature transformed phases
which is shown in [I] is not satisfied and the stretch
flange, formability deteriorates. For this reason, the
lower limit was made 5"C/sec.. By making this cooling rate
7'C/sec, it is possible to obtain stably refined l ~ w
temperature transformed phases, so the cooling race is
preferably made 7'C/sec or more,
[008.3] Further, the cooling rate in the temperature
range from 600°C to 500°C is made a relatively low speed
so as to promote the ferrite transformation which occurs
in this temperature region and make the amount of carbon
which concentrates at the austenite uniform. When the
average cooling rate in this temperature range is over
3OC/sec, the ratio of the nano hardnesses of the low
temperature transformed phases to the nano hardness of
the ferrite does not satisfy the range limited by the
above [I], so the upper 1j.mit is made 3"C/sec. if
preferably 2'C/sec, there is less fluctuation in the nano
hardnesses of the low temperature transformed phases and
less fluctuation in the yield strength.
[0084] After the sheet is cooled under the above
conditions, the sheet is run through the hot dip
galvanization bath, then is wiped by gas to adjust the
basis weight. In some cases, the sheet is then run
through an alloying furnace so as to make the base iron
diffuse in the galvanization layer for alloying
treatment. The temperature of the alloying furnace is
adjusted by the line speed. It is sufficient to select
the temperature at which alloying is completed. The
tem.perature usually becomes 460 to 600°C in range. If
460°C or less, the alloying becomes slower and the
productivity is poor. Further, if over 600°, ferritepearlite
transformation occurs and the properties
deteriorate.
[ O Q 8 5 ] After that, skin pass rolling and tension
leveling and other shape correction steps are performed
to complete the product. The draft of the skin pass
rolling is preferably 0.1 to 1.5% in range. If less than
0.1%, the effect is small and the control also is
difficult, so this becomes the lower limit. If over 1.5%,
the productivity remarkably falls, so this is made the
upper limit. The skin pass niay be performed in-line or
off-line. Further, the skin pass of the targeted draft
may be performed at one time or may be performed divided
into several times. Further, trimming etc. may be
performed.
[0086] The type of the annealing furnace used may be
any type such as the NOF-RF type or all radiant tube
furnace type. Further, it is also possible to adjust the
dew point, atmospheric ingredients, etc. for control of
the plateability. Further, it is also possible to perform
Ni or other electroplating before the continuous hot dip
galvanization line for the purpose of improving the
plateability. Further, it is also possible to perform
various types of post-treatment for imparting corrosion
resistance and other properties after the plating.
Examples
100871 Next, the present invention will be explained
in detail using invention examples.
[6888] Invention Example 1
Slabs of the symbols A to AQ which have the chemical
ingredients which are shown in Table 1 were hot rolled at
the slab heating temperature and final rolling
temperature which are shown in Table 2, water cooledat
the water cooling zone, then coiled at the temperatures
which are shown in Table 2. The final rolling
temperatures were the Ar3 point or more in each case. The
hot rolled shee-ts were pickled, then cold rolled to
obtain cold rolled sheets. The hot rolled sheet
thicknesses, cold rolling rates, and cold rolled sheet
thicknesses are shown in Table 2.
ED0893 After that, the sheets were heat treated and
hot dip galvanized by a continuous hot dip
galvannealization facility. The hot dip galvanized steel
sheets were processed by conditions shown in Table 2 such
as the average heating rate from 600°C to the annealing
temperature, the annealing temperature, the holding time,
the average cooling rate from the annealing temperature
to 650JC, and the average cooling rate from 600°C to
500°C. The sheets were run through a galvanization bath,
then cooled down to room temperature by a 1OoC/sec cooling
rate down to room temperature, then rolled by a 0.3%
draft by skin pass rolling.
[QO90] Further, cold rolled steel sheets which were
produced under the same conditions were annealed under
the same conditions, run through a galvanization bath,
then run through an alloying furnace for alloying
treatment. The alloying treatment temperature was
selected in accordance with the line speed from a 460°C to
600°C range of tenperature. After alloying treatment, the
sheet was cooled down to room temperature by a 1O0C/sec
cooling rate down to room temperature, then rolled by a
0.3% draft by skin pass rolling. The basis weight was
made about 50 g/mZ at the two sides.
[00911 Each of the obtained hot dipped galvanized
steel sheets was subjected -to a tensile test and measured
for the YS (yield strength), TS (tensile strength), and
El (elongation). Note that, the yield strength was
measured by the 0.2% offset method. The tensile test was
performed by obtaining a JIS No. 5 test piece from a 1.4
mrn thiclc sheet in a direction perpendicular to the
rolling sheet and evaluating it for tensile properties.
From the measurement values, the following properties
were evaluated. The results are shown in Table 7.
[0092g [Tensile Strength (Stress) (TS) I
The case of a tensile strength of 780 MPa or more is
skiown as "G" (good) and the case of less than 780 MPa is
shown as "P" (poor) in Table 2.
[00931 [Yield Ratio]
The ca'se of a yield ratio of 0.7 or less is shown as "G"
(good) and -the case of over 0.7 is shown as "P" (poor) in
Table 2. 0.7 or more is sufficient.
[0094] [Strength-Ductility Balance]
The strength-ductility balance (TSxEl[MPa.%]) was found
and used as an indicator of the press formability. The
results are shown in Table 2. The symbols are shown
below. 14000 or more is sufficient.
VG (very good): 16000 or more,
G (good): 15000 to less than 16000,
F (fair) : 14000 to less than 15000,
P (poor) : less than 14000.
[00!35l [Fluctuations in Yield Strength]
Next, to evaluate the fluctuations in the yield strength,
the annealing temperature was changed and hot dipped
galvanized steel sheets and hot dipped galvannealed steel
sheets were produced. That is, slabs which were cast by
the same ingredients were used to prepare sheets under
the same hot rolling conditions and cold rolling
conditions. At the annealing step, the heating and
cooling conditions were made the same but the annealing
temperatures were changed in the range of 720 to 800°C.
The sheets were run through a plating bath, then cooled
down to room temperature by a 10°C/sec cooling rate down
to room temperature, then rolled by a 0.3% draft by skin
pass rolling or were alloyed, then cooled down to room
temperature by a lO0C/sec cooling rate down to room
temperature, then rolled by a 0.3% draft by skin pass
rolling for the test. The tensile,characteristics of
these steel sheets were evaluated. At that time, two or
more points were measured in each of the ranges of 720 to
73OoC, 730 to 740°C, 740 to 750°C, 750°C to 760°C, 760°C to
770°C, 770°C to 780°C, 780°C to 790°C, and 790°C to 800°C.
Paeferably, measurement data is obtained for three or
more points. In the present embodiment, a plurality of
coils were used for measurement while changing the
anneal'ing temperatures. A single coil may also be used
for measurement while changing the annealing
temperatures. The symbols of the fluctuations in the
yield strength are shown below. 60 MPa or less i.s
sufficient.
VG (~rery good): Difference of maximum value and minimum
value of yield streng-th of 40 MPa or less when making
- range of annealing temperature 120 to 80O0C,
G (good): Difference of maximum value and minimum value
of yield strength of over 40 MPa when making range of
annealing temperature 720 to 800°C,
P (poor): Difference of maximum value and minimum value
of yield strength of over 60 MPa when making range of
annealing temperature 720 to 800°C,
The results are shown in Table 2.
[0096] [Average Grain Size of Low Temperature
Transformed Phase,?]
The microstructure was examined by polishing the sheet
thickness cross-section, then using the LePera method for
corrosion and observing the cross-section by a power of
lO00X by a metal microscope. With LePera corrosion, soft
ferrite phases are colored and hard low temperature
transformed phases remain white. Due to this, the
fraction of the ferrite phase and the average grain size
of the low temperature transformed phases were round. The
average grain size was determined by using a grid of
length 1.5 pm squares and using the point count to find
the area of the low temperature transformed phases. The
structural photograph and grid are' 'shown in FIG. 2. The
number of counted points was 200 points. Further, the
number of low temperature transformed phases which were
contained in the region for finding the area rake by the
point cdunt was counted and the average diameter when
making the crystal grains circular were - calculated from
the area rate and the number. The results are shown in
Table 2.
E00971, [Strength-Hole Expandability Balance]
The strength-hole expandability balance was evaluated by
the stretch flange formability. The stretch flange
formability was evaluated using the hole expandability
value h by the hole expandability test which is shown in
the Japan Iron arid Steel Federa-tinn Standard JFSTi001-
1996. In this case as well, the strength-hole
expandability value balance (TSxh[MPa~%j) was found and
used as an indicator of the stretch flange formability.
The results are shown in Table 2. The symbols are shown
below. 20000 or more is sufficient.
VG (very good) : 24000 or more,
G (good): 22000 to less than 24000,
F (fair): 20000 to less than 22000,
P (poor): less than 20000.
[0098] [Fluctuations in Nano Hardness]
The nano hardness was measured by the ultrasmall loading
hardness method p~escribed in JIS % 2255. The measured
load was made 1 FN. The average nano hardness was
measured at 100 points. For both the ferrite haidness and
the low temperature transformed phase hardness, the steel
sheet was cut, the sheet thickness cross-section was
polished, then the cross-section was electrolytically
corroded so bring out the microstructure, an SPM image
was observed, and the ferrite phase and the low
tempera-ture transformed phases were judged arid the nano
hardness measured. The fluctuations in the nano hardness
of the low temperature transformed phases are judged by
the ratio of -the low temperature transformed phases
contained in the range of 1 to 5 tin:es the average
hardness of the ferrite phase. The results are shown in
Table 2. Thc symbols are shown below. 80% or mvii. is
sufficient.
VG (very good) : loo%,
d (good): 90% to less than loo%,
F (fair): 80% to less than 90%,
P (poor): less than 80%.
[cJcJS9] [Spot Weldability]
The spot weldability was evaluated as follows: Electrode
(dome type) : tip diameter 6 mm4, pressiiig force: 4.31tN,
welding current: current (CE)kA right before occurrence
of spatter and (CE+1.5)kA, wel-ding time: 15 cycles,
holding time: 10 cycles. After welding, a cross tensile
test was performed in accordance with JIS Z 3137. Welding
was performed 10 times by a welding current of ICE) la.
The lowest value among those was defined as CTSICE). As
opposed to this, the lowest value of the CTS when welding
10 times by a welding current of the spatter region of
(CEfl. 5) kA was defined as CTS (CE+l. 5) . The ratio of these
values (=CTS (CE+l. 5) /CTS (CE) ) was used to evaluate the
weldability as follows. 0.7 or more is sufficient.
G (good) : 0.8 or more
F (fair) : 0.7 to less than 0.8
P (poor) : less than 0.7
~ o ~ o o % [Plateability] [Alloying Reactivity]
The plateability and alloying reactivity were evaluated
as follows: 'The symbols which show the plateability are
shown below.
G (good) : No non-plating,
F (fair) : Some non-plating,
P (poor) : Much non-plating.
The symbols which show the alloying reactivity are shown
below.
G (good): No uneven alloying in surface appearance,
F (fair): Some uneven alloying in surface appearance,
P (poor): Much uneven alloying in,surface appearance.
The above results are shown in Table 2. There is no
problem so long as not "P".
gO%O%] F'trom Table 2 which summarizes the resillts, it
is learned that the steel sheets of the present invention
are excellent in all of the steel sheet shapeability,
~Eldability, and plateability and have little
fluctuations in material quality.
[0102Q Invention Example 2
Slabs of the symbols AR to BA which have the chemical
ingredients which are shown in Table 1 were processed
under the production conditions which are shown in Table
2 in the same way as in Invention Example 1 to pfoduce
hot dipped galvannealed steel sheets. Note that the
finishing temperature at the time of ho-t rolling was the
Ar3 point or more in each case. The inventors used this
experiment to study the effects of addition of Sn on
alloying of the galvanization.
[0103] As the method of evaluation, using a continuous
hot dip galvanization line, sheets were run under the
annealing conditions which are shown in Table 2, then
were run through a galvanization bath. The lowest
temperature at which alloying sufficiently occurred was
found from Experiment Nos. 44 and 49 to which Sn was not
added. After that, Experiment Nos. 45 to 48 and 50 to 53
were run to search for the lowest temperature at which
alloying sufficiently occurred. The difference from the
lowest temperature which was found by Experiment Nos. 44
and 49 was found.
[01041 The results are shown together in Table 2. Due
to this, it was learned that alloying was promoted by the
addition of Sn. However, in Experiment Nos. 48 and 53
where the amounts of Sn were made more than the limit,
defects occurred in the hot rolled sheet. Further, the
additional evaluation which was performed in Invention
Example 1 was also performed and the results are shown in
Table 2. Due to these, it is learned that the steel
sheets of the present invention are excellent in all of
the shapeability, weldability, and plateability and are
small in fluctuation of material quality.
[0105] Invention Example 3
Slabs of the symbols A, C, and H which have the chemical
ingredients which are shown in Table 1 were processed
uEder the production conditions which are shown in Table
2 in the same way as in Invention Example 1 to produce
hot dipped galvanized steel sheets and hot dipped
galvannealed steel sheets. Note that the finishing
temperature at the time of hot rolling was the Ar3 point
or more in each case. The inventors used this experiment
to study the production conditions of the steel sheets.
The results of evaluation are shown in Table 2. Due to
this, it is learned that the steel sheets of the present
invention are excellent in all of the shapeability,
w e l d a b i l i - t y , and p l a t e a b l l i t y and a r e small i n
f l u c t u a t i o n of m a t e r i a l q u a l i t y .
eologj

[0107% Table 2
/ I 1 I 1 1 Hot 1 ,_,, 1 1st 1 2nd 1 3rd 1 nth 1 5th 1 . / Heatlng 1 I 1 I Cooling I
S t e e l S1.ab Flnlsh- L U A U pass pass pass pass pass LUAU
Colllng rate from , . I I , , 1 r a t e from1 /Exp./Steel/ type h e a t i n g ' i i g ,..Ir,-",";:,"P.lro 1jingl cold cold / cold I cold cold r?lled 1 600°C to AC3 Hold=annngea lins!

Slab
heating
temp.
("CI
1250
1250
1250
1250
1250
1250
1250
1250
1250
1250
1250
5th
pass
cold
rolling
rate
1 % )
15.0
15.0
10.0
10.0
10.0
15.0 1
1
15.0
10.0 1
15.0
10.0 1
10.0
15.0 1
10.0
10.0
10.0
10.0
15.0
10.0
10.0
10.0
steel
type
class
(Table 1)
1nv.e~.
1nv.e~.
1nir.e~.
1nv.e~.
1nv.e~.
Inu.ex.
1nv.e~.
Exp.
no.
58
59
60
61
62
63
64
Steel
type
C
1 C
C
C
C
C
C
Finishing
temp.
('CI
900
900
900
900
900
900
900
900
900
900
900
cold
sheet
thickness
(-I
1.4
1.4
0.8
1.4
1.4
, l.4
1.4
1.4
1. 4
1.4
1.4
1.4
1.4
1.4
1.4
1 . 4
1.4
1.4
1.4
1.4
1.4
-1.4
1.4
1.4
1.4
1.4
1.4
1.4
1.4
69
70
71
72
65
66
67
68
e m .
(TI
780
780
780
780
780
780
780
780
710
720
745
780
795
830
862
---1-- -
780
' 780
780
780
780
780
780
780
780
720
780
780
, 780
900
900
900
900
900
900
900
900
900
900
900
900
900
900
900
900
900
900
Heating
r6a0t0e0 Cf rom
annealing
temp.
(.C/S)
1.9
1.9
3.9
0.3
-------1.9 ---
3.5
5.2
7.1
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
1.9
7
1.9
1 . 9
1.9
1.9
1.9
1.9
tine
120
120
60
120
120
, 120
120
120
120
120
120
120
120
120
120
5
---
25
120
300
120
120
120 pppp
120
120
120
120
120
120
120
temp.
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
843
i 843
843
843
843
843
843
843
843
843
coiling
(OC1
550
550
550
550
550
550 I
550
550 1
550
550 1
550 1
C
C
C
/ C
C
C
C
C
550 1
550
550 1
550 1
550 1
550
550
550 1
550 1
i 550 1
550
/
550
550
550 '
550
550
550
Cooling 1
rate from
annealing1 I
to 650% 1
(OC/5j ;
11.0 /
11.0 '
19.0 1
1 11.0 I
11.0 1
A,
1nv.e~. I 1250
1nv.e~. 1 1250
1nv.e~./ 1250
1nv.e~./ 1250
/ 1nv.e~.
Inv.ex.
1nv.e~.
1nv.e~.
Hot
rS",y:p
thickness
(rn)
2.4
4.0
4.0
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
11.0
11.0
11.0
11.0
1.1.0
11.0
11.0
11.0
11.0
11.0
11.0
11.0
11.0
3.2
5.8
8.2
11. .0
23.0
-73 1 C
74 / C
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
2.8
,
0
1
1
75
76
77
78
79
80
81
82
83
84
85
86
Inv.ex./12501
1nv.e~. / 1250
1nv.e~./ 1250
"ld
rolling
rate
(a)
41.0
65.0
80.0
50.0
50.0
1 53.0
50.0
50.0
50.0 1
50.0 1
50.0 1
11.0 1
11.3 1
11.0 1
11.0
11.0 _
C
C
C
C
C
C
C
C
C
C
C
C
1nv.e~.
1nv.e~.
1iiv.e~.
inv.ex.
Inv.ex.
Inv.er.
1nv.e~.
1nv.e~.
1nv.e~.
1nv.e~.
1nv.e:~.
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
50.0
1250
1250
i250
1250
12501
1250
1250
1250
1250
1250
1.250
1st
pass
cold
rolling
rate
(%]
17.0
21.0
44.0
15.0
15.0
20.0
15.0
15.0
20.0
15.0
15.0 ,
15.0
15.0
15.0
15.0
~ ~ p p - - ~ ~ - - ~ 50.~0 ~ ~ - 20.0
20.0
15.0
50.0
20.0
20.0
23.0
15.0
50.0
15.0
20.0
15.0
20.0
4th
pass
cold
rolling
rate
( % I
18.5 1
1
10.0
11.0
10.5 I
11.0
10.0 1
14.0 1
10.0 1
11.0 1
2nd
pass
cold
rolling
rate
( % )
ppppp--------------
16.9
20.0
43.0
13.5
15.0
12.8 ,
15.0
13.5 1
15.0 1
13.5 1
15.0 1
10.0
11.0 1
11.0 1
10.0 1
,
14.0 1
10.5 1
11.0
10.5
14.0
10.5
11.0
10.0
10.5
11.0
10.5
3rd
pass
cold
rolling
rate
1 % )
14.5
20.0
37.5
11.0
13.5
11.0
13.5
11.0
14.5
11.0
13.5
13.5
15.0 1
15.0
13.5
15.0
12.8
15.0
12.8
15.0
12.8
------
15.0
13.5
12.8
15.0
12.8
11.0
13.5
13.5
11.0
14.5
11.0
13.5
11.0
14.5
11.0
13.5
11.0
11.0
13.5
11.0
Exp.
no.
Steel
type tyipl rate
Steel
type
class
(Table 1) rate
Slab
heating
temp.
loci rate rate
----- --(mi - 1 % ) ( % I 1%) ($1 1%) ---(.C/S, I
87 H Ilv. x. 1150 880 8 20.0 150 14.5 14.0
Finishing
temp.
(OCI rate
103
104
temp. to 650°C I
H
H
Coiling
tempt
"')
1st
pass
cold
rolling
. .
1nv.e~.
1nv.e~.
105
106
107
Hot
rolled
sheet
thickness
50.0
50.0
50.0
Cold
rolling
rate
1150
1150
H
H
H
2nd
pass
cold
rolling
20.0
50.0
15.0
108
109
110
111
112
113
11.4
115
4th
pass
cold
rolling
3rd
pass
Cold
roiling
880
880
12.8
1.0
15.0
13.5
12.8
1nv.e~.
1nv.e~.
1nv.e~.
12.8
15.0
ii
H
H
H
H
H
H
H
p5atshs
cold
rolling
550
550
11.0
13.5
14.5
11.0
11.0
1150
1150
1150
11.0
13.5
1nv.e~.
1nv.e~.
1nv.e~.
1nv.e~.
1nv.e~.
1iiv.e~.
1nv.e~.
'Old
rolled
sheet
thickness
2.8
2.8
10.5
11.0
14.0
10.0
10.5
880
880
880
10.5
11.0 J
1000
1150
1150
1150
1150
1350
1150
50.0
50.0
10.0
10.0
15.0
10.0
1nv.e~./ 1150
550
550
550
Heating
rate from
6000C
annealing
10.0
10.0
880
880
880
880
880
880
880
2.8
---2-.8 --
2.8
temp.
Annealing
t(eemCp).
20.0
15.0
1 . 4
1.4
1.4
1.4
1.4
1.4
1.4
880 1.4
1.4
1.4
1 . 4
550
550
550
550
550
550
550
"lding
time
15.0
---13.5 --
1.9
1.9
1.9
1.9
1.9
1.9
1.9
550 1.9
Cooling ~
rate from!
anneailng 1
temper. I
1.9
~ -1.9 - ~ 1.9
2.8
2.8
2.8
2.8
2.8
2.8
2.8
14.5
11.0
780
780
780
780
780
780
780
2.8 780
862
7eO
180 ~~~~~~~
50.0
50.0
50.0
50.0
50.0
50.0
50.0
14.0
10.0
835
835
835
835
835
835
835
835
20.0
50.0
50.0
150
20.0
15.0
20.0
50.0
835
835
835
50.0
15.0
120
300
120 -
120
120
120
120
120
11.0 !
11.0 I
3.2 I
1 5 . 8
8.2
11.0 I
23.0 1
11.0 1
120
5
25
1.4
1.4
11.0--1
11.0 i 1
11.0 1
1.9
1.9
795
830
835
835
120
120
. ' I
11.0
11.0 1 h
1 1 1 I 1 1 1 0 1 . 1 1st 1 2nd 1 3rd 1 4th I 5th / ,_ ,, / Heat ing / 1 I / Coolug I
Exp.
no.
Steel
type
Steel
type
class
(Table 11
Slab
heating
temp.
(*C!
Finishing
temp.
("Cl
Coiling
temp).
(°C! :ze ness
immi
L"LU
railing
rate
t$pl
pass pass
cold cold
rolling rolling
rate rate
i i l I , e l
pass
cold
rolling
rate
1 % )
pass
cold
rolling
rate
1%)
pass
cold
rolling
rate
181
L U A U
rO1le*
sheet
thickness
(mml
rate from
6000C to
annealing
temp.
~ ~ C I S I
Annealing
temp.
(OC!
Ac3
temp. Ho:tm2 rate from
annealing
temper.
to 650-C
1 0 r / q i
Exp.
no.
149 R 1nv.e~. 1250 940 570 2.8 50.0 15.0 15.0 13.5 11.0 10.0 1.4 1.9 780 843 120 11.0 I 1507- -
1250 0 570
---
2.8 5C.0 20.0 ' 11.0 10.5 10.5' 1.4 1.9 780 843 120 11.0 -- ~~~~~
---
151 R 1nv.e~. 1250 900 570 2.8 1 30.0 14.0 10.0 9.5 2.0 1.9 780 843 120 11.0
152 R 1nv.e~. 1250 900 570 2.8 1 50.0 15.0 15.0 13.5 11.0 10.0 1.4 1.9 780 843 120 11.0 I
5th
rolling
rate
(%i
10.0
10.0
15.0
Steel
type
"Id
rolled
sheet
thickness
1.4
1.4
1.4
1.4
Steel
type
class
(Table 1)
145
1-?6)
147
148
Heating
rate from
600DC to
annealing
temp.
- - (*C/S) - - 1.9
1.9
1.9
1.9
Slab
heating
temp.
l°C1
570
570
570
570
R
R
R
R
, 20.0
15.0
15.0
50.0
Annealing
temp.
(OCl
780
780
780
780
ing
temp.
l"C1
2.8
---
2.8
2.8
2.8
1nv.e~.
1nv.c~.
1nv.e~.
1nv.e~.
12.8
15.0
13.5
Ac3
temp.
843
843
843
843
50.0
50.0
50.0
50.0
1250
1250
i250
1250
11.0
~ ~ ~13.5 ~ ~ 11.0
HO:::eng
("
120
120
120
120
Hot
Finish-Coiling rolled
rolling
1
900
900
900
1st
pass
cold
rolling
rate
(%i
10.5
11.0
10.0
Cooling
rate from
annealing
temper.
to 650°C
(OC:si
8.2 ,
11.0
23.0 -
11.0 I
(mm) 1 % )
2nd
pass
cold
rolling
rate
( % )
3rd
rolling
rate
1%)
4th
rolling
rate
1%)
Strengthhole
expandability
balance
ratio
1
2
3
4
5
-6
7 1
8 1
9
-10 1
11 1
12
13
14
15
16
. 17
18
19
20
21
22
Pot
weld-
Ratio of
nano
hardness
Tensile
strength
Fluc. of
nano
hardness
a-grain
Of
Sheet
(pm)
cooling
rate
from
600°C t o
5000C
!.C/S)
Exp.
no.
A
B
C
D
E
F
G
H
I
J
K
L
N
N
--0 -
P
Q
R
S
T
U
V
Plating
wettabilityability
Strengthductility
balance
steel
type
Ferrite
phase
fraction
( % )
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
23
24
25
Average
grain
size of
lob temp.
transformed
phases
!pml
G
G
G
G
G
G
Alloying
16
16
16
16
16
16
14
14
13
1 13
12
11
1 6
16
16
16
18
18
18
18
17
17
W
i
Y
Remarks
Pluc.
of
yield
stress
G
G
G
G
G
G
G
Class
I
! I
7 3
72
7 3
7 4
7 1
7 3
--76 -
7 7
77
7 6
7 3
77
85
8 6
-88
86
68
7 5
72
7 3
7 3
64
0.7
0.7
0.7
26 Z 0.7
0.7
0.7
0.7
-
27
28
29
1.13 1
0.45 1
0.52
0.52 1
0.57 1
0.70 1
1.32
0.42 1
0.67
0.66
0.67
VG
VG
1- VG
---VG --
VG
VG
VG
PA
AB
AC
VG
VG
VG
VG
VG
VG
VG
12 / 81
1.2
2.3
2.1
2.2
2.1
2.3
1.2
2.8
2.2
2.3
2.3
VG
VG
VG
VG
VG
VG
VG
0.58
0.59
0.59
12
12
12
12
12
12
0.57 -
0.34
0.44
0.34
83
------.-.--- 83
---82 --
8 1
8 3
82
P
G
VG
VG
G
G
-P
F
VG
F
F
1nv. ex.
Comp ex
Inv. ex.
Comp. ex.
1 Inv, ex.
0.72
1.13
0.52
0.55
1
0.41
0.67
0.61
0.58
0.57
G
G
-G
P
G
G
P
2.2
2.1
2.2
F
P
G
VG
G
P
F
VG
G
-P -
F
1 1nv. ex.
1 camp . ex.
1.9
0.8
1.9
2.0
1.9
4.2
2.8
2.3
2.2
2.1
G
VG
VG
2.2
2.2
2.1
2.2
P
G
G
G
I G
G
P
G
G
G
G
1.19
G
G
G
G
G
G
G
A G
G
VG
G
G
P
G
G I
G
G
G
G
G
G
G
G
F
-G
-G
G
G
G
P
G
G
G
G
G
P
G
G
- - - - . . . . . . . - - ~ - - -G ~ ~ ~ ~
G
G
P
G
G
G
P
G
G
G
G , r
P
G 1
VG
VG
VG
VG 1
--P
F
VG
VG
VG
1 VG
P
VG
~ VG~ - ~ G
P
G
VG
VG
VG
P
G
G
G
VG
P
VG
VG
P
P
G
VG
VG
VG
P
G
G
G
G
G
G
G
G
G
G
G
G
G
G -
G
G
G
G
G r
G
G
P
VG , G
VG 1 G
G i~
G I P
I P I G
G
G
G
G
G
G
G
G
G
P
G
G
G
G
G
G
G
G
G
Comp. ex.
G
VG
VG
P
P
G
G
G
G
G
G
G
G
G
G
G
G
G
P
G
G
G
VG
VG
G
P
G
VG
VG
G
G
VG
VG
VG
VG
VG
VG
G
G
VG
G
G
-
G
Inn"v.. e x '"i
~nve.x.
Comp. ex
Comp. ex.
I
1nv.e~.
ID". ex.
Comp. ex.
In". ex.
I
G
Comp. ex. I
Comp. ex.
Conp. ex.
Inv. ex,
Inv. ex.
Comp-ez-;
Comp. ex.
111". EX.
Inv. ex.
In". ex.
Inv. ex.
,
Pot
weld-
G
G
G
EX^.
no.
-30
3 1
32 1
steel
type
AD
Ferrite
phase
fraction
( $ 1
61
83
82
Plating
wettabilityability
G
2
G
Ratio of
nano
hardness
2.9
-----2.5 -1
2.3
33
34 1
35 /
Average
grain
size of
l o b temp.
transphases
( P i
0.55
0.55
0.54
G I G 1
G 1
G 1
-
G 1
G 1
G 1
G 1
G 1
G I G
G 1
G I
G 1
G
G
G
G ,
G
G
G
G -
Cooling
rate
from
600°C to
500°C
("C/s)
0.7
Alloying
G
- G
G
Fluc. of
nano
hardness
G
G
VG
8 1
83
64
a-grain
size of
sheet
(pm)
. 16
G
G
- G
G
G
G
G
G
G
G
G
GGI
G
G
--G
G
G
G
G
G
G
G
-G
AG 1
AH 1
A1 1
AE 16
57
58
2.2 1
1.8 1
2.3 1
0.57
1.37
0.55
Fluc.
of
yield
stress
P
G
VG
---
AF 1
r-m
37 /
38 /
39 1
40
41 /
42 /
-43 /
44 1
-45 I
-46 /
47 1
48
-49
50 1
51 1
52 1
53
-
54
55
56
G
G
G
VG
VG
P 1
-P
P 1
P
P 1
P 1
P 1
P
- P 1
0.7
0.7
1.0
VG
VG
VG
VG
VG
VG
VG
----V-G -
VG
-VG
VG
VG
VG
VG
VG
VG
G
G
VG
P
VG
C
C
G
G
G
G
Remarks
-
16
16
I4
0.7
- G I G ]
G G ~
G 1
G 1
G 1
G 1
G 1
G
G
G
G
G
G
G
G
G
G
G
G
G
AJ I
AK 1
AL 1
P 1
AN
A0 1
AP /
AQ 1
AR 1
AS 1
AT 1
AU
AV
An1
AX
AY
AZ
BA
c
C
C
G
VG
G VG
z","ii,"h
G
G
G
Class
Comp. ex.
Inv. ex.
Inv--e-x.
-comp.ex.
16
1.0
1.0
Comp . ex.
1nv.e~.
Inv. ex.
G
G
G
G
G
1nv.e~.
Comp. ex.
Comp . ex.
-
Comp . ex.
Comp . ex.
Comp . ex.
Comp . ex.
Comp . ex. I
Co:np. ex.
Comp.ex.
~nve.x.
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
-1.0
1.0
1.0
1.0
1.0
/ 1.0
1.0
1.0
-
~~~~~~
1.0
1.0
1.0
-G 1
G 1
G 1
Defects
in hot
Strengthductility
balance
--------P -------
A
VG
G 1
G 1
G
G
G
G
G
G
G
G
P
G
Strengthhole
expandability
balance
VG
- VG
VG
14 64 0.48 2.3 G G P
23
18
G
G
14
14
14
14
14
14
14
14
14
2
14
14
14
14
1.40
0.78
sheet
Defects
in hot
sheet
72
7 3
G
P
P
-----
65
64
65
64
8 6
8 4
86
8 4
---83 --
84
8 4
83
83
, 84
2.3
2.1
1nv.e~.
Inv. ex.
1nv.e~
1nv.e~.
C0mp.e~.
comp. ex.
Comp.e >:. 1
G
P
VG
I
G
- G
0.49
0.52
0.54
0.54
0.53
0.52
0.52
0.57
0.57
0.58
0.59
0.59
0.55
0.55
G
G
14 1 83 0.53
0.52
0.55
0.72
0.52
0.52
14
14
2 6
/ 22
16
2.2
2.2
2.3
2.1
2.3 '
2.1
2.2
2.3
2.3
2.1
2.3
2.2
2.1
2.1
G
G
82
83
8 6
8 4
73
1.9
2.0
2.0
. 3
i.1
2.1
VG
G
VG
G
VG
G
--------G ----
F
-F F
F
T
VG
VG
G
G
VG
G
G
G
G
VG
VG
VG
-
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
P
P
G
G
G
G
G
G
G
G
G
G
G
G
G
~~--G --------~~~~
G
7
P
P
P
VG
VG
VG
VG
VG
VG
----------ppp
VG
VG
VG
VG
G
G
G
P
P
G
VG
, VG
VG
~ ~ ~ - - - - - P
------P
VG
E ~ ~ .
no.
59
60
61
62
63
64
Plating
wettabilityability
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
c
Alloying
G
G
G
G
G
-G G
G
G
G
G
G
G
G
G
G
G
G
G
G
G 8
G
G
G
G
G
G
G
c
Ferrite
phase
fraction
( % I
73
72
73
73
72
73
a-grain
Size Of
Sheet
(PI
12
7
2 1
16
15
15
steel
type
C
C 1
C
C
C
C
F~uc.
of
yield
stress
VG
P
VG
VG
VG
G
VG
VG
G
VG
VG
dG-.,-VGp pp
G
VG
P
VG
VG
VG
P ppppp
G
VG
VG
VG
P
G
VG
VG
VG
c
Coaling
rate
from
600°C to
500%
( O c / s )
1.0
1.6
1.0
1.0
1.0
1.0
--P
F
F
G
VG
VG
VG
P
VG
VG
VG
VG
VG
VG
a.
P
F
G
VG
VG
F
VG
P
VG
VG
VG
P
F
65 / C 1 1.0 15 7 8 1.12 2.1 G P
Average
grain
size of
lo* temp.
transformed
phases
(pa)
0.43
1.72
1.52
0.52
0.52
0.52
G
G
L--,,-pppp G
G
G
G
G
G
G
G
G
G r
66
67
68
69
70
71
72
73
74
75
76
77
Remarks
---
1
Pitting
Deterior-
"y,","epf
Shape
G
ppppp
G
G
G
G
G
G
; G
79
Ratio of
nano
hardness
1.9
1.8
1.7
L . A
- .
2.i
2.1
Class
Inv.ex.
Comp. ex.
Comp. ex.
Inv. ex.
1nv.e~.
1nv.e~.
Comp . ex.
Comp. ex
Inv.-e-l- :.
Inv.ex.
1nv.e~.
Inv. ex.
In\-.ex.
~amp.ex.
Comp. ex.
Inv.e -x.
1nv.e~.
Inv. ex.
Comp.ex.
Inv.e x-.
In-i .ex.
Inv. ex.
Ina.e x.-
Comp. ex.
1nv. 6x7
1nv.e~.
1nv.e~.
Inv-ex,
C
C
C
C
C
C
C
C
--C
C
C
C
78
C
Fluc. of
nano
hardness
VG
VG
G
VG
VG 1
G 1
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
C
1.0
80 / C
12 I
13
14
16
-pppp 17
19
22
11
-1 3
16
2 0
16
I n 7 7 R n n nn i n P P
1.0
1.0
3.8
2.5
1.4
I. . 0
0.5
81
82
83
84
85
86
1.0
16
C
C
C
C
C
C
Strengthhole
expandability
balance
VG
P
P
Strengthductility
VG
G
VG
VG
VG
G
ratio
G
G
G
G
G
G
95
8 9
8 8
7 5
73
7 1
61
8 3
7 8
75
7 3
7 4
16
16
16
16
16
16
16
Pot
weld-
G
G
G
strength
G
P
G 1
G 1
G
G !
16
~ 73 ~ ~ - ~ 0.08
0.62 ~ ~ p ~ ~ p ~ - 0.52
0.52
0.52
0.73
0.73
0.08
0.52
0.52
0.52
1.80
7 4
73
7 4
73
75
73
7 4
VG
VG
F
73
0.62
i G
G
G
3.6
2.5
2.1
2.1
2.1
1.8
4.0
2.1
2 . 1.
2.1
2.1
0.52
0.42
0.55
0.55
0.53
0.57
0.61
0.76
2.0
G 1
G 1
G 1
VG 1
VG
F
P
VG
VG
VG .
P
1.9
1.8
2.1
2.2
1.9
2.2
- .
2.3
G
P
G
G
G
G
G
P
G
G
G
G
VG
VG
P
G
G
VG
VG
P
G
G
G 1
G 1
1.8mGG Gpp
G
G
G
G
G ,.
G
G VG
G G
VG
VG
VG
P
F
UG
VG
VG
G
G
G
G
G
G
G
G
G
G
G
G
~ ~ ~ G
G
VG
JIG
VG
VG
VG
VG
VG
Strengthductility
balance
ppppppp
P
VG
VG
VG
G 1
VG
VG
VG
G
P
F
F
G
VG
VG
VG
P
VG
VG
VG
VG
VG
VG
VG
VG
VG
VG
VG
Ratio of
hardness
-1.3
2.1
2.3
/.I
" .
1.9
1.7
-1.7 -
2.1
2.1
2.1
2.1
2.4
2.2
2.0
2.2
1.9
1.9
3.8
2.2
2.2
2.2
2.2
2.1
2.2
2.1
2.1
&." - .
2.0
~xp.
no.
-
-88
89
90
91
92
93
Strengthhole
expandability
balance
G
VG
P
Ferrite
phase
fraction
1 % )
83
73
74
75
72
Fluc. of
nano
hardness
F
-----
G
F
F
F
7 -
F
F
-F
G
F
F
F
F
F
F
P
F
pp--p-
F
F
P
F
F
F
ppp---
F
P
F
Average
grain
size of
temp.
transformed
phases
( P I
0.63
901
1.50
0.66
0.41
1.22
Steel
type
H
H
H
H 1
H 1
H
Pot
weldability
G
~~-ppp-pppp G
G
1.44
0.43
0.44
0.45
1.07
0.09
0.62
0.49
0.50
0.49
0.64
0.71
0.09 1 0.49
0.48
1.27
1.60
0.88
0.62
0.43
1.00
0.52
0.49
-9 4 '
95
96
97
98
'-991
100
101
102
103
1.04
105
106
-107
108
109
110
111
112
r- 1.13
114
115
116
Yield
ratio
P
G
G
G
G
G
G
G
G
P
G
G
G
G
G
G
P
G
G
G
G
G
G
G
G
G
G
Cooling
rate
from
600°C to
5000C
(-C/S)
1.0
1.0
1.0
1.0
Plating
wettability
G
G
G
strength
P
G
G 1
G 1
G 1
P I
G I
G
G
G
P
P
G
G
G
G
G
G
G
G
----G
G
G
G
G
G
G
G
G
H '
H
H
H
H
H 1
H
H 1
H
H
H
H
ti
H
H
H
H
H
H
H
H
H
H
a-grain
of
sheet
(I""'
---
--. 23 -
16
24
13 VG
VG
P
P
VG
VG
F
P
ppp-pp
P
F
G
VG
VG
F
VG
P
VG
VG
P
P
F
VG
A
VG
P
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
G
--G -
G
G
G
G
G
G
G
G
G
1.0 / 1.3
1.6 / 7
G
G
G
G
G
G
G
G
G p---------p
G
G
G
G
G
G
G
1 G
G
G
G
G
G
-G
G
G
Alloying
G
G
G
F 1 G
1.0 / 21 / 73
G
G
G
G
G
G
G
G 1
G 1
G 1
G 1
G 1
G 1
G
G
G
G !
G
G
G
G 1
-G
G
- G
G
F~uc.
of
yield
stress
-
G
G
G
7 4
75
77
93
87
8 4
7 1
7 4
7 3
63
8 8
83
75
74
75
74
/ 74
75
7 4
75
7 4
1.0
1.0 / 16
1.0 / 1 6
1.0 / 15
1.0 / 12
1.0 I 13
1.0 / 14
1.0 / 16
1 . / 17
G
G
P
-VG
G
G
G
VG
VG
G
G
G
G
G
VG
P
G
G
P
P
G
G
G
P
G
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
1.0
3.8
2.5
Remarks Class
I
' 2 0
22
12
--13
------16-
2 0
17
17
17
-16
1.6
16
16
1
I
1
1
1
-1
1
1
1
1
1
Pitring 1
Deterioration
of
sheet
shape
-
1
1
1
-1
1
i
-
Comp. ex
Inv. ex.
camp. ex-
1nv.e~.
1nv.e~.
Comp . ex
-camp. ex
1nv.e~.
1nv. ex.
inv. ex.
~ompe.x
Camp. ex
In7.e~.
1nv.e~.
Inv.ex.
1nv.e~.
Inv.ex.
Comp . ex
Camp. ex
Inv.ex.
1nv-.- e~.
comp . ex
Comp . ex
1nv.e-~ .
In". ex.
1n.r.e~.
1nv.e~.
camp;;
inv. ex
I -
Average
Cooling grain
rate a-orain
Ferrite size of Strengthleu"
temp, Ratio of Fluc. of Strength- hole Pot Plating F~uc.
E ~st~ee.l from
Size Of phase
no. type 600°C to
fraction trans- nano nano ductility expand- weld- wett- Alloying yield Remarks Class
500nC sheet formed hardness hardness
ratio strength balance ability abilityability
iF6i balance Stress
i0c/si
(prni phases
ipm) ~~~-~~~
117 H 1.4 16 7 5 0.48 1.9 F G G VG VG G G G T G Inv. ex
~ ~ ~ ~ ~ ~ 118 H / 3.0 16 7 1 0.44 2.3 FIG G I VG F G G G G Inv. ex
~~~~~~~
119 H 0.5 16 7 4 0.61 3.0 F G G VG VG G G G G Inv. ex
120 R 1.0 / 26 87 0.75 A.. L- F P P P G I G G G G Comp . e:
121 R 1.0 / 22 85 0.82 1.0 F P P P G G G G G Cornp . e:
122 R 1.0 / 16 75 0.64 2.3 F G G VG ] VG G G G G ID". ex
123 / R 1.0 1 26 2 1.50 2.2 G G G aP ICJ G G G Comp. e:
'
-- -----
124 R 73 0.81 2.1 F G G VG VG G G G G Inv. ex
125 R 1.0 13 7 4 0.44 2.0 F G G VG VG G G G G Inv. ex
126 R 1.6 8 75 1.49 1.6 F G P G P G G G P Comp. e,
127 R 1.0 23 75 1.55 1 1.6 F G G VG P G G G VG Comp. e;
128 R 1.0 18 74 1 0.54 1 2.3 P/G -G VG -VG G G -G G -Inv. ex ~ ~~~~~~
129 R 1.0 17 75 0.55 2.1 F G G VG VG G G G G In7,. ex
130 R 1.0 17 7 4 0.54 2.3 F G G G F G G G G Inv. ex
131 R 1.0 17 73 1.26 1 2.2 G P P P G G G VG Comp. e, -
132 R I 1.0 13 94 0.11 3.7 G P P F P G G G VG Comp . e?
133 R I 1.0 I. 4 85 0.61 2.4 F G G F F G G G G 1 n v . e ~
134 R I 1.0 15 8 4 0.55 2.2 F G G VG VG G G G G 1nv.e~
135 F ( 1 0 18 7 4 0.51 2.2 F G G VG VG G G G G Inv. ex
136 R I 1.0 19 73 0.52 2.3 F. G G VG VG G G G G Inv. ex - p p 137 R / 1.0 2 1 7 1 0.82 1.7 F G G VG F G G G G 1nv.e~
143 R I 1.0 1 18 1 74 1 1.90 1 2.3 1 P I G / G 1 VG 1 P G I G 1 G [ P I 1 ~ompe.,
8144 R 1.0 18 73 -& , - F G G VG F I G G G I G I 1 n v . e ~ : -
145 R 1 1.0 1 18 / 74 1 0.66 1 2.1 1 F G I G I V G / V G G I G / G 1 G 1 1 1nv. ex
1 1 1 1 I 1 Average
Coollng gram
rate a-waln
S1Ze Of
Ferrlte sr$e of
EX^. Steel from phase low temp
600DC to annealed
no. type fxact~on trans-
500nC (%) formed
lpm) phases
Ratio of
nano
hardness
Strength
flu^. of Strength- hole
nano
Tens'1e ductlllty expandratlo
strength
hardness balance ab~lxLy
balance
Pot Plating Fluc.
weld- wett- Alloying Of erna arks yield
bility ability stress
Class
-
[-nv.ex.
[-nv.er.
-a.m p . ex
[nv.ex.
[-nv.-e x.
[nv.ex.
[-nv.e x.
Industrial Applicability
[0108] The present invention stably and inexpensively
provides steel sheets with a high strength of a tensile
strength of 780 MPa and an excellent shapeability which
are suitable for chassis parts which are used for an
automobile. It promises to greatly contribute to the
lighter weight of au-tomobiles and is extremely high in
effect in industry.

Claim 1.
High strength hot dipped galvanized steel sheet
characterized by containing, as ingredients of the steel,
by mass%,
C: 0.05 to 0.1%,
Si: 0.1 to 1.0%,
Mn: 2.0% to 2.5%,
Al: 0.02 to 0.1%,
Ti: 0.01 to 0.05%,
Cr: 0.1 to 1.0%,
Sn: 0.0010 to 0.1%, and
a balance of Fe and unavoidable impurities,
having a microstructure comprised of low temperature
transformed phases of a ferrite phase fraction of 70 to
90% and a balance of martensite,
having an average grain size of the low temperature
transformed phases of 0.1 to 1 pm,
having a ratio of average nano hardnesses of the
ferrite phase and the low temperature transformed phases
of 1.5 to 3.0, and
having a nano hardness of the low temperature
transformed phases at 80% or more oE the measurement
points of 1 to 5 times the average nano hardness of the
ferrite phase.
Claim 2
A method of production of high strength hot dipped
galvanized steel sheet characterized by heating a slab
which has the steel ingredients as set forth in claim 1
to 1000 to 1350°C, then hot rolling at a final rolling
temperature Ar3 or more, coiling at 600°C or less,
pickling, cold rolling at a rolling rate of 30 to 70%,
and, after that, heat treating while making a temperature
of 720°C to a temperature of ths lower of 850°C or the Ac3
temperature the annealing temperature, during which
heating in the temperature range from at least 600°C to
the annealing temperature by a 0.5OC/sec to b°C/sec
heating rate, holding at the annealing temperature for 10
see or more, then cooling in at least the temperature
range of the annealing temperature to 650°C by a cooling
rate of 5'C/sec or more, further cooling in at least the
temperature range of 600°C to 500°C by a coollng race of
3'C/sec or less; then performing hot dip galvanization or
hot dip galvanneallzation.