We claim:
1. A high-strength hot-dip galvannealed steel sheet excellent in bake hardenability,
comprising:
a base steel sheet containing, in mass,
C: 0.075 to 0.400,
Si: 0.01 to 2.00,
Mn: 0.80 to 3.50,
P: 0.0001 to 0.100,
S: 0.0001 to 0.0100,
Al: 0.001 to 2.00,
N: 0.0001 to 0.0100,
O: 0.0001 to 0.0100 each,
with the balance made up of Fe and inevitable impurities,
wherein a structure of the base steel sheet contains, in volume fraction, 3 or more of
a retained austenite phase, 50 or less of a ferrite phase, and 40 or more of a hard
phase,
at a range from 1/8 thickness centered around a 1/4 sheet thickness from a surface to
3/8 thickness centered around the 1/4 sheet thickness from the surface at the base steel
sheet,
an average dislocation density is 5  1013/m2
or more,
a solid-solution C amount contained in the retained austenite phase is in mass 0.70 to
1.00,
an X-ray random intensity ratio of FCC iron in an texture of the retained austenite phase
is 3.0 or less,
a ratio between a grain diameter relative to a rolling direction and a grain diameter
relative to a sheet width direction of the retained austenite phase is 0.75 to 1.33,
further, a hot-dip galvanized layer is formed at the surface of the base steel sheet and
the sheet thickness becomes 0.6 to 5.0 mm.
2. The high-strength hot-dip galvannealed steel sheet excellent in the bake hardenability as
claimed in claim 1, wherein the hard phase is made up of one or both of a bainitic ferrite
phase and a bainite phase, a tempered martensite phase, and a fresh martensite phase. 3. The high-strength hot-dip galvannealed steel sheet excellent in the bake hardenability as
claimed in claim 1,
wherein oxides are finely dispersed, and a decarburized layer whose thickness is 0.01
to 10.0 m is formed at a surface layer portion of the base steel sheet, and an average
grain diameter of the oxides is 500 nm or less, an average density is 1.0  1012 oxides/m2
or more.
4. The high-strength hot-dip galvannealed steel sheet excellent in the bake hardenability as
claimed in claim 1, wherein the base steel sheet containing, in mass, one kind or two
kinds or more from among
Ti: 0.001 to 0.150,
Nb: 0.001 to 0.100,
V: 0.001 to 0.300.
5. The high-strength hot-dip galvannealed steel sheet excellent in the bake hardenability as
claimed in claim 1, wherein the base steel sheet containing, in mass, one kind or two
kinds or more from among
Mo: 0.01 to 2.00,
W: 0.01 to 2.00,
Cr: 0.01 to 2.00,
Ni: 0.01 to 2.00,
Cu: 0.01 to 2.00,
B: 0.0001 to 0.0100.
6. The high-strength hot-dip g galvannealed steel sheet excellent in the bake hardenability
as claimed in claim 1, wherein the base steel sheet containing, in mass, one kind or
two kinds or more from among
Ca, Ce, Mg, Zr, La, REM for 0.0001 to 0.0100 as a total.
7. A manufacturing method of a high-strength hot-dip galvannealed steel sheet excellent in
bake hardenability, wherein an alloying treatment is performed for the hot-dip galvanized layer is formed at the surface of the high-strength steel sheet according to claim 1, said
method comprises:
a hot-rolling step of heating a slab having a chemical component containing, in mass,
C: 0.075 to 0.400,
Si: 0.01 to 2.00,
Mn: 0.80 to 3.50,
P: 0.0001 to 0.100,
S: 0.0001 to 0.0100,
Al: 0.001 to 2.00,
N: 0.0001 to 0.0100,
O: 0.0001 to 0.0100 each,
with the balance made up of Fe and inevitable impurities to 1180C or more, starting a
hot-rolling performed by plural passes, performing the hot-rolling in which a
relationship among a temperature “T” of a hot-rolled steel sheet within a range from
1050C to a rolling completion temperature, a sheet thickness h, and an elapsed time t
between each pass satisfy the following expression (1), and completing the rolling at a
temperature range of 880C or more;
a first cooling step of starting cooling after an elapsed time after the hot-rolling
completion to the cooling start is set to be 1.0 second or more, and stopping the cooling
at 450C or more;
a cold-rolling step of setting an elapsed time after the first cooling until 400C to be 1.0
hour or more, and thereafter, a cold-rolling is performed while setting a total reduction
ratio to be 30 to 75;
a continuous annealing step of annealing at a maximum heating temperature of (Ac3 
50)C or more; and
a plating step of forming a hot-dip galvanized layer at a surface of the steel sheet by
immersing the steel sheet into galvanizing bath after the continuous annealing step,
wherein a bainite transformation process retaining the steel sheet at a temperature
within a range of 300 to 470C for 20 to 1000 seconds before the steel sheet is immersed
into the galvanizing bath or after the immersion.
[Numerical Expression 1] Here, in the expression (1), “N” represents a total number of passes from the hot-rolling
start to completion, “i” represents an order of each pass, “Ti” represents a rolling
temperature (C) at the i-th pass, “hi” represents a sheet thickness (mm) after the
processing of the i-th pass, “ti” represents an elapsed time from the i-th pass to the next
pass. Note that when i = 1, h0 = a slab thickness. Besides, an elapsed time from a final
pass to the next pass is an elapsed time from the final pass to the cooling start time after
the hot-rolling completion.
7. The manufacturing method of the high-strength hot-dip galvannealed steel sheet
excellent in the bake hardenability as claimed in claim 7,
wherein in the plating step, oxides are generated at a surface layer portion of the steel
sheet at a preheating zone where an air ratio being a ratio between a volume of air
contained in mixed gas in a unit volume in the mixed gas of air and fuel gas used for
heating and a volume of air theoretically required to enable complete combustion of the
fuel gas contained in the mixed gas in a unit volume is set at 0.7 to 1.2, subsequently
the oxides are reduced at a reduction zone where a partial pressure ratio (P(H2O)/P(H2))
between H2O and H2 is set to be 0.0001 to 2.0, and thereafter, the steel sheet is immersed
into the galvanizing bath under a condition in which a plating bath temperature is 450
to 470C, a steel sheet temperature when entering into the plating bath is 430 to 490C,
an effective Al amount in the plating bath is 0.01 to 0.18 mass, to thereby form the
hot-dip galvanized layer at the surface of the steel sheet.
9. The manufacturing method of the high-strength hot-dip galvannealed steel sheet
excellent in the bake hardenability as claimed in claim 7, wherein
the plating step is followed by a temper rolling step of performing a rolling for the steel
sheet with a reduction ratio of 5.00 or less.
10. The manufacturing method of a high-strength hot-dip galvannealed steel sheet excellent
in the bake hardenability as claimed in claim 7, wherein
the plating step and the bainite transformation process is followed by alloying the hot
dip galvanized layer.
11. The manufacturing method of the high-strength hot-dip galvannealed steel sheet
excellent in the bake hardenability as claimed in claim 8, wherein
a temper rolling step of performing a rolling for the steel sheet with a reduction ratio of
less than 10 after the hot-dip galvanized layer is alloyed.

TECHNICAL FIELD
[0001] The present invention relates to a high-strength hot-dip galvanized .steel sheet, a high-strength alloyed hot-dip galvanized steel sheet excellent in bake hardenability, and a manufacturing method thereof, BACKGROUND ART
[0002] In recent years, requirement for high-strengthening of a steel sheet used for a vehicle and so on becomes high, and a. high-strength steel sheet whose tensile maximum stress is 900 MPa or more comes to be used. On the other hand, it is required to show excellent forming workability at a forming work time such a.s'presswork, However, improvement in.,the strength easy to incur deterioration in the forming workability, and it is difficult to satisfy both requirements,
Accordingly, in recent years, a steel sheet using bake hardenability (BO-ability) by a coating/baking process (bake hardening process) after the forming work has been developed so that workability such as ductility and hole expandability and high-strengthening are both enabled,

[0003] Here, the bake hardening is a phenomenon in which C (solid-solution C) and N (solid-solution W) remaining in a steel sheet in a solid-solution state diffuse to dislocations during a baking process (normally heated up to approximately 170°C, then kept for several dozen minutes) after coating, the dislocations are fixed, and thereby, a yield strength increases.. The increased amount of.the yield strength is a coated bake hardening amount (BH amount), and the BH amount is generally known to increase by increasing a solid-solution C amount or a solid-solution N amount,
[0004] In Patent Literature 1, a cold-rolled steel sheet is disclosed in which a hard structure made up of bainite and martensite is a main structure thereof, and a high bake hardening amount is secured by limiting a fraction of ferrite into 5% or less, [0005] Besides, in Patent Literature. 2, a high-strength cold-rolled 3teel sheet is disclosed in which bainite is a main structure thereof, ..a hardness ratio between bainite and ferrite is made small, and dispersion of hardness in each structure is made small to thereby improve the bake haidenabi1ity, the ductility, and the hole expandability. [0006] Besides, in Patent Literature 3, a method is disclosed in which a steel sheet is made to be one containing tempered martensite and/or tempered bainite obtained by performing annealing for a hot-rolled steel sheet without performing- cold-rolling,

or performing the annealing for two times after the
cold-rolling to thereby improve the workability and
the bake hardenability,
[0007] Besides, in each of Patent Literatures A, b
and 6, an art improving the bake hardenability by
adding a lot of N is disclosed.
CITATION LIST
PATENT LITERATURE
[0008]' Patent Literature 1: Japanese Laid-open
Patent Publication No, 2008-144233
Patent Literature 2: Japanese Laid-open Patent Publication No, 2004-2 6 3270
Patent Literature 3: Japanese Laid-open Patent Publication No, 2003-277 88 4
Patent Literature 4: Japanese Laid-open Patent Publication No, 2005-023348
Patent Literature 5: Japanese Laid-open Patent Publication No, 2003-049242
Patent Literature 6: Japanese Laid-open Patent Publication No, .200 1 -2 47 94 6 SUMMARY OF INVENTION TECHNICAL PROBLEM
[0009] However, in. Patent Literatures 1 and 2f it is mentioned about the bake hardening amount, but anisotropy of the bake hardenability is not mentioned, and it is extremely unstable whether or not it is possible to stably secure a desired bake hardening amount,
Besides, in the method of Patent Literature 3,

the cold-rolling is not performed, and therefore, there is a problem in which sheet thickness accuracy of the steel sheet, deteriorates. Besides, even if the cold-rolling is performed, the annealing step after the cold-rolling is performed for two times, and there is a. problem in which manufacturing cost increases.
Besides, in Patent Literatures 4, 5 and 8, it is necessary to add a lot of N to secure the bake hardening amount, and there is a possibility in. which we.ldabil.ity is deteriorated, [0010] The present invention is made in consideration of the circumstances as stated above, and an object thereof is to provide a high-strength hot-dip galvanized steel sheet, a high-strength alloyed hot-dip galvanized steel sheet securing high-strength of a tensile maximum strength of 900 MPa or more, excellent ductility, and excellent in bake hardenabi1ity, and a manufacturing method thereof. SOLUTION TO PROBLEM
[0011] The present inventors studied hard to solve the above-stated problems. As a result, they found that it is possible to obtain a steel sheet whose bake hardening amount is large and having isotropic bake hardenability while securing high-strength of tensile maximum strength of 900 MPa or more, and excellent ductility by increasing an average dislocation density in the steel sheet, weakening anisotropy of a texture of -austenite, and enabling

anisotropic structure,
[0012] Summary of the present invention with the aim of solving the above-stated problems is as fallows. [0013] [1] A high-strength hot-dip galvanized steel sheet excellent in bake hardenability, includes a base steel sheet containing, in raass%, C: 0.075 to 0.400%, Si: 0.01 to 2.00%, Mn: 0.80 to 3.50%, P: 0.00O1 to 0.100%, S: 0.0001 to 0.0100%f Al: 0.001 to 2.00%, N: 0,0001 to 0.0100, Oi 0.0001 to 0.0100 each, with the balance made up of Fe and inevitable impurities, wherein a structure of the base steel sheet contains, in volume fraction, 3% or more of a retained austenite phase, 50%. or less of a ferrite phase, and 40% or more of a hard phase, at a range from 1/8 thickness centered around a 1/4 sheet thickness from a surface to 3/8 thickness centered around the 1/4 sheet thickness from the surface at. the base steel sheet, an average dislocation density is 5 x 1013/m2 or more, a solid-solution C amount ■contained in the retained austenite phase is.in mass%, 0.70 to 1.00%, an X-ray random intensity ratio of FCC .iron in an texture ox the retained austenite phase is 3,0 or less, a ratio between a grain diameter relative to a roiling direction and a grain diameter relative to a sheet width direction of the retained austenite phase is 0.75 to 1.33, further, a hot-dip galvanized layer is formed at the surface of the base steel sheet and the sheet thickness becomes 0.6 to 5.0 mm.

[2] The high-strength hot-dip galvanized steel sheet according to [1], wherein the hard phase is made up of a balnitic ferrite phase and/or a bninite phase, a tempered martensite phase, and a fresh martensite phase.
[3] The high-strength hot-dip galvanized steel sheet excellent in the bake ha.rdenabili.ty according to [1], wherein, oxides are finely dispersed, and a decarburxzed layer v/hose thickness., is 0.01 ^m to 10,0 j.im is formed at a surface layer portion of the base ■steel sheet, and an average grain diameter of the oxides is 500 nra or less, and an average density is 1,0 x 1012 oxides/m2 or more,
14] The high-strength hot-dip galvanized steel sheet excellent in the bake hardenability according to [1], further containing, in inass%, one kind or two kinds or more from among Ti: 0.001 to 0.15 0%, Kb: 0.0 01 to 0.100%, V: 0.001 to 0.300%.
[5] The high-strength hot-dip galvanized steel sheet excellent in the bake hardenability according to [1], further containing, in mass%, one kind or two kinds or more from among Mo: 0.01. to 2.00%, W: 0.01 to 2.00%, Cr: 0.01 to 2,00%, Ni: 0.01 to 2.00%, Cu: 0.01 to 2.00%, B: 0.0001 to 0.0100%.
[6] The high-strength hot-dip galvanized steel sheet excellent in the bake hardenability according to [1] , further containing, in masi;%, one kind or two kinds or more from, among Ca, Ce, Mg, Zr, La, REM for 0.0001 to 0.0100% as a total.

[7] A high-strength alloyed hot-dip galvanized steel sheet excellent in bake hardenability, wherein an alloying treatment is performed for the hot-dip galvanized layer formed at a surface of the high-strength steel sheet according to [1],
[0014] [8] A manufacturing method of a high-strength hot-dip galvanized steel sheet excellent in bake hardenabillty, including? a hot-rolling step of heating a slab having a chemical component containing, in raass%, C: 0.075 to 0.400%r si: 0.01 to 2,0 0%, Mn: 0.80 to 3.50%, P: 0.0001 to 0,100%, S: 0,0001 to 0,0100%, Al : 0.0 01 to 2,0 0%, 33; 0.0 0 01 to 0,0100, 0: 0,0001 to 0.0100 each, with the balance made up of l*e and inevitable impurities to 1180°C or more, starting a hot-rolling performing by plural passes, performed the hot-rolling in which a relationship among a temperature "T" of a hot-rolled steel sheet within a range from 1050°C to a rolling completion temperature, a sheet thickness uh", and an elapsed time "t" between, each pass satisfy the following .expression (1), and completing the rolling at a temperature range of 880°C or more; a first cooling step of starting cooling after an elapsed time after the hot-rolling completion to the cooling start is set to be 1,0 second or more, and stopping the cooling at 450°C or more; a cold-rolling step of setting an elapsed time after the first cooling until 400°C to be 1.0 hour or more, and thereafter, a cold-rolling is performed while setting a total reduction ratio to be

30 to 75%; a continuous annealing step of annealing at a maximum heating temperature of (AC3 - 50)°C or more; and a plating step of forming a hot-dip galvanized layer at a surface of the steel sheet by immersing the steel, sheet into galvanizing bath after the continuous annealing step. [0015] ^Numerical Expression 1]

o.io<

£{5.20*10B -r' -iMxitr3 -r,1 +i.«*io-r, -$.67

•,^frH-^M"',i"

CD
Here, in the expression (1) , VN" represents a total number of passes from the hot-rolling start to completion, "i" represents an order of passes, "IV represents a rolling temperature (°C) at the i-th pass, "hi" represents a sheet thickness (mm) after the processing of the i-th pass, "ti" represents an elapsed time from the i-th pass to the next pass. Note that when i =? 1, h0 = a slab thickness. Besides, an elapsed time from a final pass to the next pass is an elapsed time from the final pass to the cooling start time after the hot-rolling completion. [0016] [9] The manufacturing method of the high-strength hot-dip galvanized steel sheet excellent in the bake hardenability according to [8], wherein in the plating step, oxides are generated at a surface layer portion of the steel sheet at a preheating zone where an air ratio being a ratio between a volume of

air contained in mixed gas in a unit volume in the mixed gas of air and fuel gas used for heating and a volume of air theoretically required to enable complete combustion of the fuel gas contained in the mixed gas in a unit volume is set at 0."/ to 1.2, subsequently the oxides are reduced at a reduction zone where a partial pressure ratio (P (H20)/P (H?J ) between Hs0 and H2 is set to be 0,0001 to 2.0r and thereafter, the steel sheet is immersed into the galvanizing bath under a condition in which a plating bath temperature is 450 to 470°C, a steel sheet temperature at the time entering into the plating bath is 430 to 4 90°C, an effective A.1 amount in the plating bath is 0.01 to 0.18 raass%, to thereby form the hot-dip galvanized layer at the surface of the steel sheet.
[10] The manufacturing method of the high-strength hot-dip galvanized steel sheet excellent in the bake hardenability according to [8], further including; a temper roll.i.ng step of., performing .a. rolling for the steel sheet with a reduction ratio of 5.00% or less after the plating step.
[11] A manufacturing method of a high-strength alloyed hot-dip galvanized steel sheet excellent in bake hardenability, including: alloying the hot-dip galvanized layer after the high-strength hot-dip galvanized steel sheet is manufactured by the manufacturing method according to [8].
[12] The manufacturing method of the high-

.strength alloyed hot-dip galvanized steel sheet excellent in the bake hardenability according to [10] [10], further including: a temper rolling step of performing a rolling for the .steel sheet, with a reduction ratio of less than 10% after the hot-dip galvanized layer is alloyed. ADVANTAGEOUS EFFECTS OF INVENTION
[0017] It is possible for a high-strength hot-dip galvanized steel sheet, a high-strength alloyed hot-dip galvanized steel sheet according to the present invention to adhere C to a lot of dislocation and to increase a C amount in a steel sheet because enough average dislocation density is supplied by defining a microstructure of the steel sheet, into a predetermined fraction. As a result, it is possible to enlarge a bake hardening amount. Besides, it is possible to make retained austenite unstable by reducing a solid-solution C amount in the retained austenite and to easily transform the retained austenite into martensite by a forming work, and . s.o on, As a result, it is possible to enlarge the bake hardening amount, Further, the retained austenite transforms into extremely hard martensite as stated above, and thereby, mobile dislocation is introduced at a periphery of a martensite structure, and therefore, it is possible to secure a further bake hardening amount..
[0018] Besides, it is possible to enable an isotropic retained austenite structure by reducing an

~r pre iron of the
■ ,DT,«itv ratio oi K^L- -<-
X-ray random accordingly/ XT-
texture of the r ■ ^ the isotropic
mak„ it transform i"1-
is possible to make _ . subseqUent
, t.he processing such as tn
martensita oy the pr pos3ibie to obtain
rl therefore/ i1- -L& ^
forming work, a" „_,,„, a pattern ot
Be
v. vo hardenability. »*=-BJ-
;sotropic bake haro^ defined'
, ftf the retained austenxte i- «
crystal grams o_ - martensite
hv a transformation ratio -
and thereby, at. ..rertion ia able to be
„n a processing direct-,
depending on P hardenability
n„t1nt, and the D^V.
^ ^tablV con^i-auu/ made
made £.T.,.
is siatw" nfi the hign
^ „*lvanized steel ^eeu
ctrenqth hot-dip gaivani
strcng^- „-,ivanazed steel ^te.
v. 11ftved hot-dip gaivaxi
strength alloyed enouqh secure
1rt the present invention to
according .to tne y , the
d ,- J to obtain L,"1--
* ,ja«ina amount, ^USJ . w i ^
the bake hardening ^ ^ thej.eby possible
, .,v,roPic bake hardenability- ■ ' ,naddition
"oLropX bake hardenability in add..
to largely improve ductility.
4. in the strength ana
to the improvement , to the present
■riflS the steel sheet according
Besides, tn« ic bake
iw the isotropic "
** able to supply the
invention is abi possible to
K-ntv and therefore, it is P
hardenability, a of the
u t- limiting a usage direct
design without 1^- ■ ■ mention when
steel sheet according to the p
=, member and so on.
,nr.lied for a menu-"--*-
it is appli^-a ..]rinq method of the
. , , in the manufacturing
[00191 Besides, _ ^ ^^ and the
. - M-rength hot-dip galvanise ■ ^
high-strenyi- wvnized steel onect.
K i n h - S1
, bnt-diP qalvaniicu «-i-h al loved hoi. UJ-P J -strength an^i

sible to
.j'.-..-, i-rt f-he present inve according to int. y-
ntion, it is poss
having sufficient
en
able the desired mxoi.^u
„„^c.<■■ c development, of nd to suppress acve >. F
dislocation density, ai
■^ 4-n wpaken the anisotropy of the texture of.austenite to weaken
M. .rrurture by defining conditions of Lhc the austenite structuit 4^
u„omiBflt cooling step. ■, ■! ■ ,« «tr>n and the subsequent hot-rolling step <=nu
■4- ■«. no —ible to make the martensite Accordingly, it is po^ibie
i h« the subsequent processing j- - +- V- T n «; formed »y r. ne .juw^^-n structure transronu^w j
„ +. v,o hake hardenability • 4-v.nn-ir and to improve the riatce .isotropic aiiu
, ., a^n before and after the Besides, in the plating -ep, ■
• rqion of the steel sheet into the plating bath, immersion or i-"«
j ,4. a f-PiriDerature range or 4^oi ohfet is retained at a tempera the steel srn-«-^ ^
4. ~-i+-0 is easy to transform result, the retained austtnu
into martens
step to increase the ba
DESCRIPTION OF EMBODIMENTS
f00?0] Hereinafter,
, <. * hiah-strength alloyed hot-galvanized steel sheet, a high
1 „nf nurpl lent in bales Ivanized steer biict-t.
^ ,^ processing being a subsequent ■ ite at the proce^iuy
a nigh-strength hot-dip
dip ga
A -, manufacturing method thereof ability, and a maniucn-i-m ^
harden
according to the prese
detail. [0021]
re described in
nt invention a
Ivanized steel
trength hot-dip ga
ike hardening.amount■

The high-strength hot-dip galvanized steel sheet according to the present invention is characterized in that a base steel sheet contains, in mass%, Ci 0.075% to 0.400%, Si: 0.10 to 2.00%, Mn: 0.80 to 3.50%, P: 0.0001 to 0.100%, St 0.0001 to 0.0100%, Al: 0.001 to 2.00%, N: 0.0001 to 0.0100, 0. 0.0001 to 0.0100, with the balance made up of Fe and inevitable impurities, a structure of the base steel sheet contains, ..in volume fraction, 3% or more of a retained austenite phase, 50% or less of a ferrite phase, and 40% or more of a hard phase, at a range from 1/8 thickness centered around a 1/4 sheet
, ■ , -Fv,-,m i «„rf?ce to 1/8 thickness centered thickness from a t-urr^oe LU J/ «
around the 1/4 sheet thickness from the surface at the base steel sheet, an average dislocation density is 5 x 1013/m* or more, a solid-solution C amount contained in the retained austenite phase is in mass% 0.70 to 1.00%, an X-ray random intensity ratio of FCC
. ,-,-. ^-P t-hP rptained austenite phase is iron of a texture ol tne renaxn^u QUOV.
3.0 or. less, a. .ratio between .a grain diameter .. ....
re. au:
relative to a rolling direction and a grain diameter dative to a sheet width direction of the retained .stenite phase is 0.75 to 1.33, further, a hot-dip
galvanized layer is formed at the surface of the base
steel sheet and the sheet thickness becomes 0.6 to
5 . 0 mm.
Hereinafter, limitation reasons of a steel sheet structure and a chemical component (composition) of the present invention are described. Note that a

notation of "%" represents mass% unless otherwise
spec Ifled.
[0022] (Sheet Thickness)
A sheet thickness of a steel sheet to be applied is 0.6 to 5.0 mm, When it is less than 0.6 mm., it is not suitable because it is difficult to keep a shape of the steel sheet flat, and when it exceeds 5.0 mm, a predetermined microstructure cannot be obtained because it becomes difficult to uniformly cool inside the steel sheet,
(Microstructure)
The microstructure of the base steel sheet of the high-strength hot-dip galvanized steel sheet of the present invention has a predetermined chemical component, and contains, in volume fraction, 3% ox more of a retained austenite phase (hereinafter, referred to as retained austenite), 50% or less of a ferrite phase (hereinafter, referred to as ferrite), and 40% or more of a hard phase, at a range from 1/8 thickness centered around a 1/4 sheet thickness from a surface to 3/8 thickness centered around the 1/4 sheet thickness from the surface at the steel sheet. [0023] "Ferrite"
The ferrite is a structure whose yield stress is low and having an excellent work hardening property. Accordingly, when a ferrite fraction is excessively increased, strength before a bake hardening process increases and the yield stress after the bake hardening process decreases, and therefore, the bake

hardenab.il ity largely deteriorates. Therefore, the ferrite fraction in the steel sheet is set to be 50% or less, The ferrite fraction is preferably 45% or lass, and more preferably 40% or less to further increase the bake hardenability, A lower limit of the ferrite fraction is not particularly defined, and it may be "0" [zero}%, However, the ferrite fraction is preferably 5% or more, and more preferably 10% or more from a point of view of the ductility. [0024] "Retained Austenite"
The retained austenite is a structure having an FCC (Face-Centered Cubic lattice) crystal structure, transforming into hard martensite during processing such as a forming work, and showing large work hardening. Besides, the martensite generated daring the processing rapidly increases the yield stress thereof by being tempered at a low temperature in a bake hardening process, and therefore, a large bake hardening amount can be obtained by increasing the volume fraction, of the retained austenite.. . Further, . . the retained austenite transforms into martensite, and thereby, a mobile dislocation is introduced at a periphery of the martensite structure, and therefore, the bake hardening amount can further be obtained. From- these points of view, the volume fraction of the retained austenite is set to be 3% or more, Further, the volume fraction of the retained austenite is preferably 5% or more, and more preferably 7% or more to increase the ductility together with the bake

cannot be obtained by the retained austenite which is excessively stable for processing. When the solid-solution carbon amount in the retained austenite exceeds 1.00%, the retained austenite becomes excessively stable, a martensite amount generated by the processing such as the forming work before the bake hardening process becomes small, and the sufficient bake hardenabiltty cannot be obtained, The solid-solution carbon amount in the retained austenite is preferably 0,96% or less to efficiently transform the retained austenite into martensite. On the other hand, when the solid-solution carbon amount in the retained austenite is below 0„70%f a martensite transformation starts during a process cooling to the room temperature after an annealing step, and the fraction of the retained austenite cannot be secured, and therefore, the solid-solution carbon amount is set to be 0,70% or more. The solid-solution carbon amount i.3 preferably 0,75% or more, and more preferably 0,80% or more to obtain a sufficient amount of retained austenite, [0027] Here, both of the following affects on an adjustment of the solid-solution carbon, where 1] a rolling reduction and a temperature from 1050°C to a finish rolling completion are set to be within a range of the later-described expression {1}, and 2] as it is described later, it is kept at 300 to 470<>C for 20 to 1000 seconds after the annealing, Namely,-the adjustment of the solid-solution carbon cannot be

substantially performed unless both I] and 2] are satisfied.
[0028] Namely, when the later-described expression (1) is satisfied, the mlcrostructure of a hot-rolled sheet becomes a homogeneous and fine structure, and island pearlites disperse homogeneously and finely. In this pearlite, Mn is segregated, and therefore, it is preferentially substituted into retained y by passing through a phase transformation at the annealing step, The solid-solution carbon is efficiently concentrated according to the phase transformation at the annealing step in the homogeneous and fine austenite to be a proper solid-solution carbon amount.
i
On the other hand, when the expression {1} is below a specified range, rec.rysta.lli zat ion does not proceed, and therefore, coarse pearlite extending in a rolling direction is generated. The retained austenite generated by passing through the annealing step becomes extended coarse austenite. Accordingly, ... the concentration of carbon resulting from the phase transformation is difficult to proceed,- and the solid-solution carbon does not become a proper range. Besides, a shape of the retained austenite becomes a problem.
Besides, when the later-described expression {1) is over the specified range, the recrystallization -excessively proceeds, and massive and coarse pearlite i.s generated. The retained austenite generated by

passing through the annealing step becomes massive
and coarse austenite, Accordingly, the concentration of carbon resulting from the phase transformation is difficult to proceed, and the solid-solution carbon does not become the proper range, Besides, the shape of retained austenite becomes the problem. [002 9] Note that the solid-solution C amount (Cy) in the retained austenite is able to be found by performing an X-ray . clif fraction test under the same condition as a measurement, of an area fraction of the retained austenite, finding a lattice shape number "a" of the retained austenite, and using the following expression {2}, Note that the expression (2) is disclosed in a document "Scripta Metallurgies, et Materialia, vol.24. 1990. p509~514". [O03OJ [Numerical Expression 2]
„ (a- 0.3556) 12.01 y 0.00095 55.84
[0031] Besides, the transformation from the retained austenite into lartensite according to the processing is affected by a crystal orientation of the retained austenite. Accordingly, when the crystal orientation of the retained austenite strongly deflects, a transformation ratio into martens ite relative to a degree of processing changes depending on a processi.ng direction, and the bake hardening amount

changes. Accordingly, to obtain the isotropic bake hardening amount, it is necessary to make the crystal orientation of the retained austenite random so that the transformation ratio into martensite relative to the degree of processing is constant even if the processing is performed in any direction, [0032] As for the deflection of the crystal orientation of retained austenite, it is possible to evaluate a degree thereof by measuring a texture of the FCC crystal of iron by the X-ray diffraction method, Specifically, an X-ray random intensity ratio of the FCC iron may be found from a crystal orientation distribution function (called as an Orientation Distribution Function, ODF) representing a three-dimensional texture calculated based on plural pole figures from among {2Q0}, {311}, {220} pole figures measured by the X-ray diffraction.
In the present embodiment, it is necessary to set the X-ray random intensity ratio of the FCC iron of the texture of the retained austenite. .at .3.0. or less to sufficiently reduce anisotropy of the bake hardenabiiity, and to obtain the isotropic bake hardening amount. The lower the random intensity ratio is, the more it is preferable to reduce the anisotropy, and it is preferably 2,5 or less, and more preferably 2,0 or less. A lower limit of the random intensity ratio is not particularly limited, but it is industrially extremely difficult to set it at less than 1,2, and therefore it is preferably 1,2

or more.
Note that the X-ray random intensity ratio is a numerical value in which the x-ray intensities of each of a standard sample which does not. have an integration in a specific orientation and a sample .material are measured under the same condition by the X-ray diffraction method and so on, and the obtained X-ray intensity of the sample material is divided by the X-ray intensity of the standard sample. [0033] Manufacturing of samples for the X-ray diffraction may be performed as stated below,
At first, a steel sheet is polished to a predetermined position in a sheet thickness direction by mechanical polishing, chemical polishing, and so on, strain is removed by electrolytic, polishing and chemical polishing according to need, and at the same time, it is adjusted so that a 1/4 sheet thickness portion becomes a measuring surface. Note that it is difficult to precisely set the measuring surface at the.1/4 sheet thickness portion, and therefore, the . sample may be manufactured such that a surface within a range of 3% relative to the sheet thickness around a target position becomes the measuring surface, Besides, when the measurement by the X-ray diffraction is difficult, statistically enough, number of measurements may be performed by an EBSD method. [0034] Further, when crystal grains of the retained austenite extend in a specific direction at a surface in parallel to the sheet surface, the ' trans formation

ratio into martensite changes depending on a processing direction, and therefore, the anisotropy in the bake hardening amount is generated, Namely, a pattern of the crystal grains of the retained austenite is defined, and thereby, it is possible to weaken the anisotropy in the bake hardening amount.
In the present embodiment, when an average grain diameter of the crystal grain of the retained austenite relative to a rolling direction is set as d(RD), and an average grain diameter relative to a sheet width direction is set as d(TD), a parameter "d(RD)/d{TD)" made up of the both is limited to be a range of 0.7 5 or more and 1,33 or less, The "d (RD) /d (I'D) " is preferably 0,80 or more and 1,25 or less, and more preferably 0,85 or more and 1,18 or less to further reduce the anisotropy of the bake hardening amount,
[0035] Note that the crystal grain of the retained austenite is evaluated by finishing a surface in parallel, to the sheet surface at the 1/4 thickness, into a mirror surface, using an FE-SEM (Field Emission Scanning Electron Microscopy), and performing a high-resolution crystal orientation analysis by the H',BSD (Electron Bach-Scattering Diffraction) method, A measurement step is set as 0,1 urn, and a region in which points representing a diffraction pattern of the FCC iron aggregate 10 points or more, and a crystal disorientation with "each other is less than 10° is set to be the crystal

grain of the retained austenite. In. this case, the crystal misorientation with a second proximity measurement point is found at each measurement point, a point whose crystal misorientation is 10,0° or more which is determined to belong to a different crystal grain is ignored, and an average value of the crystal misorientations with a second proximity measurement point group whose crystal misorientation is 10,0° or less which are determined to be within the same crystal grain is found. As for the grain diameter, the grain diameters in the rolling direction and in the sheet width direction are each measured in 30 pieces to 300 pieces of retained austenite crystal grains selected at random, [0036] . "Hard Phase"
In the present embodiment, the hard phase is contained for 4 0% or more in addition to the ferrite and the retained austenite. Note that as the hard phase, it is desired that a bainitic ferrite phase, and/.or a bainite phase, a tempered martensite phase, and a fresh martensite phase are contained in complex, This hard phase is a generic of a transformation product which is harder than the ferrite phase, [0037] "Bainitic Ferrite and/or Bainite"
The- bainitic ferrite and/or the bainite is a■-structure necessary for efficiently obtaining the retained austenite, and it is preferably contained in -the steel sheet structure for 10 to 80% in volume fraction. Besides, the bainitic ferrite and/or the

bainite is a microstructure having an intermediate intensity between soft ferrite and hard martensite, the tempered martensite and the retained austenite, and it is preferably contained for 15% or more, more preferably contained for 2 0% or more from a point of view of stretch flangeability. On the other hand, when the volume fraction of the bainitic ferrite and/or the bainite exceeds 80%, it is not preferable because there is a worry that the yield stress excessively increases and the ductility deteriorates. From a point of view of the ductility, the volume fraction of the bainitic ferrite and/or the bainite
is preferably 70% or less, and more preferably 60% or 1 ess.
[0038] "Tempered Martensite"
The tempered martensite is a structure largely improving a tensile strength, and it may be contained in the steel sheet structure for 50% or less in volume fraction. From a point of view of the tensile strength, .the.volume fraction of.the tempered martensite is preferably 10% or more. On the other hand, when the volume fraction of the tempered martensite contained in the steel, sheet structure exceeds 50%, it is not preferable because the yield stress excessively increases and there is a worry that the bake hardenability deteriorates. [0039] "Fresh Martensite"
The-, fresh martensite largely improves the tensile strength, but on the other hand, it becomes a

starting point of crack to deteriorate the stretch flangebility, and therefore, it is preferably contained in the steel sheet structure for 25% or less in volume fraction. The volume fraction of the fresh martensite is preferably 20% or less, more preferably 15% or less to increase the stretch f 1 a n g e a b i 1 i t y, [0040] "Other microst matures "
Structures other than the above such as pearlite and/or coarse cementite may be contained in the microstructure of the steel sheet of the present invention, However, when an amount of the pearlite and/or the coarse cementite becomes large in the steel sheet structure of the high-strength steel sheet, the ductility deteriorates, Therefore, the volume fraction of the pearlite and/or the coarse cementite contained in the steel sheet structure is preferably 10% or less as a total, and more preferably 5% or less.
[0041] Note that, .the .. volume . fraction of each
structure contained in the steel sheet structure as described above can be measured by, for example, the method as illustrated below.
[0042] The volume fractions of the ferrite, the retained austenitey the bainitic ferrite, the ba.in.lte, the tempered martensite and the fresh martensite contained in the steel sheet structure of the -steel sheet- of the- present invention are obtained by taking--a sample of a cro-ss section in parallel to the

rolling direction of the steel sheet, and in perpendicular to the sheet surface as an observation surface, polishing the observation surface, performing nital etching, and observing a range from 1/8 thickness centered around a 1/4 sheet thickness to 3/8 thickness centered around the 1/4 sheet thickness from the surface with the field emission scanning electron microscope (FE-SEM) to measure an area fraction, [0043] (Average Dislocation Density)
The average dislocation density (before shipment) at the range of 1/8 thickness around the 1/4 sheet thickness from the surface to 3/8 thickness around the 1/4 sheet thickness from the surface of the base steel sheet according to the present embodiment is set to be 1.0 x I013/m2 or more.
Hereinafter, definition reasons of the average dislocation density are described.
It is very effective to add a lot of solid-solution carbon to increase the. bake hardenability. However, a carbon amount solid-solves in a BCC crystal of iron is very small, and therefore, it is effective that the average dislocation density in the steel sheet is increased by lowering the transformation temperature as much as possible so that carbon adheres to a lot of dislocations to increase the solid-solution carbon amount. From this point of view, the average dislocation density in the steel sheet is set to be 1.0 x 10i3/m2 or more. The

larger the dislocation density is, the easier the solid-solution carbon is obtained, and therefore, the average dislocation density is preferably 3,0 x 1013/m2 or more, and more preferably 5. 0 x 10i3/m2 or more, An upper limit of the dislocation density is not particularly provided, but it is preferably 1,0 x 1017/m2 or less, and more preferably 3.0 x 1016/m2 or less because the ductility drastically deteriorates when the dislocation density exceeds 1.0 x 1017/m7". [0044] The dislocation density is able to be found by the X-ray diffraction method and a transmission electron microscope (TEM) observation. The TEM is able to observe a minute region, and therefore, it is possible to measure each of the dislocation densities
*
of the ferrite and the hard structure in case of a multi-phase structure steel sheet. Note that in the TEM observation, it is necessary to process the steel sheet into a thin sheet state, or thin acicular at a sample manufacturing stage, and therefore, it is -di.fficu.lt to create .the., sample, and., .there is. a case when the dislocation density is lowered because the dislocation reaches a sample surface to disappear by a slight motion because the sample is small, and therefore, enough attention is required to manufacture the sample. Besides, a measurable visual field is limited in the TEM observation, On the other hand, in the X-ray diffraction method, it is possible to- relatively^ easily measure the average . dislocation density at a wide region. Accordingly,

the method measuring the dislocation density by using the X-ray diffraction method is used in the present invention,
[0045] Note that the dislocation density is obtained by controlling the fraction of the microstructure into a predetermined range, and performing an. •appropriate temper rolling, it is because the dislocation densities accumulated inside are different depending on kinds of the microstructures, [004 6] (Decarburized Layer)
Besides, in the high-strength hot-dip galvanized steel sheet of the present embodiment, the bake hardenability is improved by snaking a surface layer portion into a decarburized layer having a small amount of hard structure to disperse fine oxides to thereby increase adhesiveness of a plating layer, increase a yield stress of a base iron surface layer, and prevent to easily yield after the bake hardening process, Note that the hard structure described here is one made, up of the above-stated hard layer . and the retained austen.ite,
In the present .embodiment, a thickness of the decarburized layer formed at the surface layer portion of the base steel sheet is set to be within a range of 0.01 urn to 10,0 j.im, an average grain diameter of the oxides finely dispersed in the decarburized layer is 500 nm or less, and an average density of the oxides in the decarburized layer is within a range of 1.0 x 10'2 oxides7ma or more,

Hereinafter, limitation reasons of the above are described,
[0047] The decarburized layer having an appropriate thickness is formed at the surface layer portion of the base steel sheet, and thereby, it is possible to secure the tensile strength and to increase the adhesiveness between the base steel sheet and the plating layer. When the thickness of the decarburized layer is less than 0.01 jam, the adhesiveness with the plating layer cannot be sufficiently obtained, and therefore, the thickness of the decarburized layer is set to be 0.01 j.im or more. The thickness of the decarburized layer is preferably 0,08 ^m or more, and more preferably 0.15 Jim or more to further improve the adhesiveness with the plating layer. On the other hand, an'excessively thick, decarburized layer lowers the tensile strength and fatigue strength of the steel sheet. From this point of view, the thickness of the decarburized layer is -set to be 10,0 um or- less , From a point of- — view of the fatigue strength, the thickness of the decarburized layer is preferably 9,0 jj,m or less, and more preferably 8.0 urn or less.
Note that, the decarburized layer is a region which continues from an uppermost surface 'of the base iron, and a region whose fraction of the hard ■ structure is half or less of the fraction of the hard structure at 1/4 thickness of- the ■■■ba-se -.-steel ■■sh-e-et in the region.

The thick.nesswj.se cross section in parallel- to the rolling direction is finished into the mirror surface, it is observed by using the FE-SEM, the decarburized layer thicknesses at three points or more are measured in one steel sheet, and the average value thereof is regarded as the thickness of the decarburized layer,
[0048] A strength of the decarburized layer formed at the surface layer portion of the base steel sheet, is low, and the crack starting from the decarburized layer is difficult to occur, but there is a large strength difference between an inside of the steel sheet and the decarburized layer, and therefore, an interface between the base iron and the decarburized layer is able to function as a starting point of a new crack, To prevent the crack, it is effective to disperse the oxides into an inside of the crystal grain and/or a crystal grain boundary in the decarburized layer, the strength of the decarburized layer is. increased.to make .the strength difference with the inside of the steel sheet small, The density of the oxides is set to be 1,0 x 1012 oxides/rnz or more to obtain enough strength, The density of the oxides is preferably 3,0 x 1012 oxides/m2 or more, and more preferably 5.0 x 1Q12 oxides/m2 or more to further improve low temperature toughness. On the other hand, when the density of the oxide3 exceeds 1,0 x 1016 oxides/m2, a distance between the oxides becomes excessively near, the

surface layer portion cracks by slight processing to damage the plating layer formed thereon,- and therefore, it is set to be 1,0 x 1016 oxides/m2 or less. The density of the oxides is preferably 5,0 x 1013 oxides/m2 or less, and more preferably 1,0 x 1015 oxides/m2 or less for the steel sheet surface layer to have enough formability,
Mote that the oxide described here means the oxide mainly containing Si and/or Mn,
[0049] Besides, when a size of the oxide dispersing
in the decarfourized layer is large, the oxide in
itself functions as the starting point of crack, and
therefore, a crack resistance property such as the
ductility deteriorates, Accordingly, the average
grain diameter of the oxide is set to be 500 nrrt or
less. The average grain diameter of the oxides is
preferably 300 nm or less, and more preferably 100 nm
or less to further improve the crack resistance
property such as the ductility, h lower limit of the
average .grain diameter of the oxide is not
particularly provided, but it is necessary to
strictly control a process atmosphere and a
temperature to make it less than 30 nm, and it is
practically difficult, Therefore, it is preferably
3 0 nm or more. - ......
[0050] As for the oxide in' the decarburizud layer, the thicknesswiae cross section in parallel to the roiling direction is finished into the mirror-- surface, and observed by using the FE-SEM. The oxide density

is found by counting the number of oxides by-observing the ciecarburized layer for an extent of 7 ,um , or by using an observation area required to count up to 100Q pieces of oxides. Besides, the average grain diameter of the oxide is an average of circle-equivalent grain diameters of 100 pieces to 1000 pieces selected at random.
[0051]
Note that a hot-dip galvanized layer formed at the surface of the high-strength hot-dip galvanized steel sheet, according to the present embodiment may be performed an alloying treatment to be a high-strength alloyed hot-dip galvanized steel sheet, The high-strength alloyed hot-dip galvanized steel sheet thereby obtained is able to show the effects similar to the above-stated high-strength hot-dip galvanized steel sheet.
[0052] Besides, in the high-strength steel sheet of the present invention, a coating film made up of a phosphorus oxide and/or a composite oxide containing phosphorus may be formed at a surface of the hot-dip galvanized layer or an alloyed hot-dip galvanized layer. It can be made function as a lubricant when the steel sheet is processed, and it is possible to protect galvanization formed at the steel sheet surface,
[0053] Next,.the chemical component (composition) of-the high-strength hot-dip galvanized steel sheet and

the high-strength alloyed hot-dip galvanized steel sheet of the present invention is described. -Note that a sign [%] in the following description represents [raass%] . [005 4] "C: 0.07 5 to 0.4 0 0%"
C is contained to increase the strength and the bake hardenabiiity of the high-strength steel sheet, However, when a content of C exceeds 0.400%, the weldability becomes insufficient. From a point of view of the weldability, the content of C is preferably 0,300% or less, and more preferably 0.250% or less. On the other hand, when the content of C is less than 0.075%, the strength is lowered, and it becomes difficult to secure the tensile maximum strength of 900 MPa or more. The content of C is preferably 0,085% or'more, and more preferably 0,100% or more to further increase the strength and the bake hardenability, [0055] "Si; 0.01 to 2,00%"
-Si is an element necessary for increasing the strength and lormability by suppressing generation of iron-based carbide, and obtaining a predetermined amount of retained austenite in. the annealing step. However, when a content of Si exceeds 2,00%, there is a case when the steel sheet is embrittled, and the cold-roiling becomes difficult to perform, Accordingly, from a point of view of the cold-rolling, the content of Si is preferably 1.80% or less, and more preferably 1.50% or less. On the other hand,

when the content of Si is less than 0.01%, a lot of iron-based carbides are generated in the annealing step, the sufficient amount of retained austenites cannot be obtained, further a lot of coarse iron-based carbides are generated during the alloying treatment of the plating layer, and there is a possibility in which the strength and the formataiiity deteriorate. Accordingly, the content of Si is
preferably 0,20% or more, and more preferably 0.50% or more.
{0056] "Mn: 0.80 to 3.50%"
Mn is added to the steel sheet of the present
invention to increase the strength of the steel sheet. However, when a content of Mn exceeds 3.50%, a coarse Mn concentrated portion is generated at a sheet thickness center portion of the steel sheet, the embrittlement is easy to occur, and a trouble such, that a casted slab cracks is easy to occur. Besides, when the content of Mn exceeds 3,50%, the voidability also, deteriorates. Accordingly.,.,the content of. Mn is..,. necessary to be set at 3.50% or less. From a point of view of the voidability, the content of Mn is preferably 3.00% or less, and more preferably 2,70% or less. On the other hand, when the content of Mn is less than 0,80%, a lot of soft structures are formed during cooling after the annealing, and it becomes difficult to secure the tensile maximum strength of 900 MPa or more. Therefore, it.- is necessary to set the content of Mn at 0.80% or more.

The content of Mn is preferably 1,00% or more, and more preferably 1.30% or more to further increase the strength,
[0057] WP; 0,0001 to 0,100%"
P tends to segregate to the sheet thickness center portion of the steel sheet, and embrittles a weld zone, When a content of P exceeds 0,100%, the weld zone is drastically embrittled, and" therefore, the content of P is limited to be 0.100% or less. From a point of view of the embrittleinent, the content of P is preferably 0.030% or less. Note that the effect of the present invention is shown without particularly limiting a lower iim.it of the content of P, but manufacturing cost drastically increases if the content of P is set to be less than 0.0001%, and therefore, 0.0001% is set to be the lower limit value. Besides, it is preferably 0.00 10% or more.
[0058] WS: 0.0001 to G.0100%"
S adversely affects on the weldability and
.inan.uf acturability at the casting time... and . the,, hot-
rolling time. An upper limit value of a content of S
is therefore set to be 0.0100% or less, Besides, S
bonds to Mn to form coarse MnS and lowers the
ductility and the stretch flangeability, and
therefore, it is preferably 0.0050% or less, and more
preferably 0,0025% or less. The effect of the
present invention is shown without particularly
limi-ting a lower limit, of the content -of S-f but- ■■
manufacturing cost drastically ■increases if the

content of S is set to be leas than 0.0001%, and therefore, 0.0001% is set to be the lower limit value. Note that it is preferably 0,0005% or more, and more preferably 0,0 010% or more. [0059] WA1: 0,001% to 2,00%"
Al is an element suppressing the, generation of the iron-based carbide to make it easy to obtain the retained austenite, and increasing the strength and the formability, However, when a content of Al exceeds 2.00%, the weldability deteriorates, and therefore, an upper limit of the content of .Al is set to be 2,00%. From this point of view, the content of Al is preferably 1.50%.or less, and more preferably 1.20% or less. On the other hand, tthe effect of the present invention is shown without particularly limiting a lower limit of the content of Al, but Al is the inevitable impurity minutely existing in a raw material, and manufacturing cost drastically increases if the content of Al is set to be less than G-.001%, a-nd therefore,, the lower- limit is. set to -be ■■■■■ 0.001% or- more. Besides, Al is an effective element also as a deoxidizer, but the Al amount is preferably 0.010% or more to obtain the effect of deoxidation more sufficiently,
[0060] WN: 0,0001 to 0,0100%"
N forms a coarse nitride, and deteriorates the ductility and the stretch flangeability, and therefore, an addition amount thereof is necessary to be suppressed. When a content of N exceeds 0,0100%,

the tendency becomes obvious, and therefore, an upper limit of the N content is set to be 0.0100%. From points of view of the ductility and the stretch flangeability, the upper limit of the W content is preferably 0.0070%, and more preferably 0.0050%. Besides, N accounts for blowhole occurrence at the welding time, and therefore, the smaller the content is, the better. The effect of the present invention is* shown without particularly limiting a lower limit of the content of N, but manufacturing cost drastically increases if the content of N is set to be less than 0.0001%, and therefore, the lower limit is set to be 0.00 01% or more. Besides, it is preferably 0,0005% or more, and more preferably 0. 0 0.10% or more. [0061] "0: 0.0001 to 0,0100%"
O forms an oxide, and deteriorates the ductility
and the stretch flangeability, and therefore, an
addition amount thereof is necessary to be suppressed.
When a content of 0 exceeds 0.0100%, the ■ ■ ■ >■ ■■ ■ .
deterioration of the stretch flangeability becomes* obvious, arid therefore, an upper limit of the 0 content is set to be 0.0100% or less. Further, the content of 0 ia preferably 0.0070% or less, and more preferably 0.0050% or less. The effect of the present invention is shown without particularly limiting a lower limit of the content of 0, but ■manufacturing cost drastically increases- if the -■ ■ ■■ content of O is set to be less than 0.0001%, and

therefore, the lower limit is set to be 0.0001%, Further, the lower limit of the content of 0 is preferably 0,0003%, and more preferably 0.0005%, [0062] Besides, it is preferable to add one kind or two kinds or more from among Ti: 0.001 to 0,150%, Nb: 0,001 to 0,100%, V: 0.001 to 0,300% in addition to the above-described elements to the base steel sheet of the present invention, [0063] "Ti: 0,001 to 0,150%"
Ti is an element contributing to the strength increase of the steel sheet by precipitate strengthening, fine grain strengthening by a growth suppression of ferrite crystal grains, and dislocation strengthening through suppression of recrystaliization, However, when a content of Ti exceeds 0,150%, a precipitation amount of carbonitrides increases, and the formab.ility deteriorates, and therefore, the content of Ti is preferably 0,150% or less. The effect of the present .invention is shown without particularly limiting a lower limit of the content of Ti, but the content of Ti is preferably 0,001% or more to fully obtain the strength increasing- effect by the addition of Ti, [0064] "Mb: 0.001 to 0,100%"
Nb is an element contributing to the strength increase of the steel sheet by the precipitate strengthening, the fine grain strengthening by the growth suppression of the ferrite crystal grains, and the dislocation strengthening through suppression of

the recryatallization. However, when a content of Nb exceeds 0.150%, a precipitation amount of carbonitrides increases, and the formability deteriorates, and therefore, the content of Nb is preferably 0,150% or less. The effect of the present invention is shown without particularly limiting a lower limit of the content of Nb, but. the content of Nb is preferably 0.001% or more to sufficiently obtain the strength increasing effect by the addition of Nb, [0065] "V; 0.001 to 0.300%"
V is an element contributing to the strength increase of the steel sheet by the precipitate strengthening, the fine grain strengthening by the growth suppression of the ferrite crystal grains, and the dislocation, strengthening through suppression of the recrystallizatiori, However, when a content of V exceeds 0.300%, a precipitation amount of carbonitrides increases, and the formability deteriorates, and therefore., the content, of, V is. preferably 0.300% or less, The effect of the present invention is shown without particularly limiting a lower limit of the content of V, but the content of V is preferably 0.001% or more to sufficiently obtain the strength increasing effect by the addition of V. [0066] Further, it is preferable to add one kind or two kinds or more from among Mo: 0.01 to 2,00%, W: 0. 01 -to--2. 00%, Cr: 0.01 to 2.00%, -Ni : 0.01 to 2,00%, Cu: 0.01 to 2,00%, B; 0,0001 to 0,0100% in addition

to the above-described elements to the base steel sheet of the present invention, [0067] "Mo: 0,01 to 2, 0 0%"
Mo is an effective element suppressing a phase transformation at high temperature, and high-strengthening, Mo may be added instead of a part of C and/or Mn, When a content of Mo exceeds 2,00%, workability in hot working is damaged and productivity is lowered, and therefore, the content of Mo is preferably 2,00% or less, The effect of the present invention is shown without particularly limiting a lower limit of the content, o.f Mo, but the content of Mo is preferably 0.01% or more to sufficiently obtain the high-strengthening- by the addition of Mo, [0068] "W: 0 , 01 to 7, , 00%"
W is an effective element suppressing the phase transformation at high temperature, and high-strengthening,- m may be added instead of a part of C and/or Mn, When a content of W- exceeds .2. 00%, -the workability at hot working is damaged and productivity is lowered, and therefore, the content of W is preferably 2.00% or less. The effect of the present invention is shown without particularly limiting a lower"limit of the content of W, but the content of W is preferably 0,01% or more to sufficiently obtain the high-strengthening by the addition of- W,-[0069] "Cr: 0,01 to 2.00%"

Cr is an effective element suppressing the phase transformation at high temperature, and high-strengthening, Cr- may be added instead of a part of C and/or Mn. When a content of Cr exceeds 2.00%, the workability at hot working is damaged and the productivity is lowered, and therefore, the content of Cr is preferably 2.00% or less, The effect of the present invention is shown without particularly limiting a lower limit of the content of Cr, but the content of Cr is preferably 0,01% or more to enough obtain the high-strengthening by the addition of Cr, [0070] "Hi: 0.01 to 2.00%"
Hi is an effective element suppressing the phase transformation at high temperature, and high-strengthening, Ni may be added instead of a part of C and/or Mn, When a content of Ni exceeds 2.0 0%, the voidability is damaged, and therefore, the content of Ni is preferably 2,00% or less. The effect of the present invention is shown without particularly limiting; a lower limit of the content-.of Ni, but ..the content of Ni is preferably 0,01% or more to sufficiently obtain the high-strengthening by the additi.on of Ni. ■,
[0071] *Cui 0,01 to 2,00%"
Cu is an element increasing the strength by existing in the steel as fine particles, and it can be added instead of a part of C and/or Mn. When a content- of Cu-exceeds '2-,00%, the -voidability is damaged, and therefore, the content of Cu is

preferably 2.00% or less. The effect of the present invention is shown without particularly limiting a lower limit of the content of Cu, but the content of Cu is preferably 0.01% or more to sufficiently ofotalr the high-strengthening by the addition of Cu. [0072] WB: 0.0001 to 0.0100%"
B is an effective element suppressing the phase transformation at high temperature, and the high-strengthening, and it may be added instead of a part of C and/or Mn, When a content of B exceeds 0.0100%, the workability at hot working is damaged and the productivity is lowered, and therefore, the content of B is preferably 0.0100% or less, The effect of the present invention is shown without particularly limiting a lower limit of the content of B, but the content of B is preferably 0.0001% or more to sufficiently obtain the high-.st rengt hening by the addition of B.
[0073] Further, it is preferable to add one kind or two kinds or more from among Ca, Ce, Mg, 2r., La, REM to the base steel sheet of the present invention for 0.0001 to 0.0100% as a total in addition to the above-stated elements,
[0074] "One kind or Two kinds or more from among Ca, Ce, Mg, Zr, La, REM for 0.0001 to 0.0100% as Total"
Ca, Ce, Mg, Zr, La, REM are elements effective for improvement of the forj.nab.i...].ity, and one kind or two kinds or more can be added. However, when a total content of one kind or two kinds ox more of Ca,

Ge, Mg, Zr, La, REM exceed(s) 0.0100%, there is a possibility in which the ductility is conversely damaged, Therefore, the total content of each element is preferably 0,0100% or less, The effect of the present invention is shown without particularly limiting a lower limit of the total content of one kind or two kinds or more of Ca, Ce, Mg, Sr, La, REM, but the total content of each element is preferably 0,0001% or more to enough obtain the improving effect of the forraability of.the steel sheet.
Note that REM is an abbreviation of a Rare Earth Metal, and indicates an element belonging to a lanthanoid series, In the present invention, it is often the case that REM and Ce are added as a mi sen metal, and there is a cs.se in_which the la.nthanoi.de series elements are contained in complex in addition to La and Ce, The effect of the present invention is shown even if the lanthanoi.de series elements other than these La and Ce is contained as the inevitable impurities.
[0075] Manufacturing Method of High-Strength Hot-Dip Galvanized Steel Sheet>
Next, a manufacturing method of a high-strength hot-dip galvanized steel sheet of the present embodiment is described,
The manufacturing method of the high-strength hot-dip galvanized steel sheet of the present embodiment includes: a hot-rolling.step of- heating a ■ slab having the above-stated chemical component to

1180°C or more, starting a hot-rolling performed by plural passes, and performing the hot-rolling in which a relationship among a temperature "T" of a hot-rolled steel sheet, a sheet thickness "h", and an elapsed time "t" between each pass within a range from 1050°C to a rolling completion temperature satisfies the following expression (1), and finishing the .rolling at a. temperature range of 880°C or more; a first cooling step of starting a cooling after an elapsed time after the hot-rolling completion until the cooling start is set to be 1.0 second or more, and .stopping' the cooling at 4 5n°C or more; a cold-rolling step of setting an elapsed time after the first cooling until 400QC to be one hour or more, and thereafter, performing a cold-rolling while setting a total reduction ratio at 30% to 75%; a continuous annealing step of annealing at a maximum heating temperature (Ac3 - 50}°C or more; and a plating step of immersing the steel sheet in galvanizing bath and forming a hot-dip galvanized layer a.t a surface.of . the steel sheet after the continuous annealing step,
Note that, in the following expression (1), "N" represents a total number of passes from the hot-rolling start to completion, «j[" represents an order of each pass, UT|" represents a rolling temperature
(°C| at the i-th pass, "hi" represents a sheet thickness (mm) after the processing of the :i.-th pass, a-nd-"ti" represents an elapsed time- from-the i-th pass to the next pass.' Mote that when i = 1, ho"= a slab

thickness. Besides, an elapsed time from a final pass to the next pass is an elapsed time from the final pass to the cooling start time after the hot-rolling completion.
[0076] Here, the expression (1) is an empirical formula in consideration of a progress of recrystallization of austenite at the hot-ro 1.1 ing step and a growth of recrystallized austenite, and is an index representing a size of an austenite crystal grain after the rolling. h product of a polynomlnal ■ of the rolling temperature "T" and the reduction ratio represents a driving force of the recrystalligation. An exponential term represents tendency for dispersion of atoms, and relates to a growth rate of recrystallized grain austenite, The time "t" is added to the product of both, and thereby, a degree of growth of austenite by each one pass can be evaluated, arid an evaluation of an austenite grain diameter by a multi-pass hot-rolling can be done by finding a root mean square and-a square root thereof.
Hereinafter, limitation reasons of the above-stated manufacturing conditions are described, [0077] [Numerical Expression 3]
... (1)
[0078] To manufacture the high-strength hot-'dip

galvanized steel sheet of the present embodiment, at first, a slab having the above-described chemical component (composition) is casted.
A slab manufactured by a continuously cast slab, a thin slab caster, and so on can be used as a slab supplied for the hot-rolling, The manufacturing method of the high-strength steel sheet of the present invention conforms to a process such as a continuous casting-direct rolling (CC-DR) performing the hot-rolling just after the casting. 10079] (Hot-Rolling Step)
In the hot-rolling step, a slab heating temperature is -set to be It.l8G°C or more, When the slab heating temperature is excessively low, there is a possibility in which anisotropy of crystal orientation of a slab structure is generated resulting from the casting, Besides, when the slab heating temperature is low, a finish rolling temperature is below an Ar3 transformation point to be a two-phase, region rolling of ferrite .and. austenite, . a hot-rolled sheet structure becomes a heterogeneous mixed grain structure, the heterogeneous structure is not settled even though passing through the cold-rolling and annealing steps, and the ductility and bendability deteriorate, Besides, the lowering of the finish rolling temperature incurs excessive increase of rolling load, and there are possibilities in which- the rolling becomes difficult and -a ■ defective shape of the steel sheet after the rolling

is incurred, and therefore, the slab heating temperature is preferably 1200°C or more. On the other hand, an upper limit of the slab heating temperature is not particularly limited, but it is necessary to input a large amount of energy to heat the slab over 1300°C, and therefore, the slab heating temperature is preferably 1300°C or less. [0080] Note that the Ar3 transformation point temperature is calculated by the following expression.
Ar3 = 901 - 325 x C + 33 x Si - 92 x (Mn + Ni/2 + Cr/2 + Cu/2 + Mo/2) + 52 x Al
In the above expression, C, Si, Mn, Ni, Cr, Cu, Mo, Al represent contents [masa%] of each element. [0081] The( slab is heated up to the slab heating temperature, and thereafter, the hot-rolling is performed to make it the hot-rolled steel sheet. When the hot-rolling is performed, the texture of austenite becomes strong and the anisotropy thereof also becomes large if excessive rolling reduction is applied at high temperature, .-.To avoid this, the .hot-rolling is performed under the condition satisfying the above-stated expression (1) made up of the temperature of the hot-rolled steel sheet, the sheet thickness, and the elapsed time between each pass during a period from 1100°C to the hot-rolling complet ion,
The above-stated expression (1) is also an expression evaluating a -degree of- development- of t-h-e texture of austenite, and when a value of the above-

stated expression (1) is below 0,10, the texture of austanite at the hot-rolled steel sheet becomes strong, and therefore, the value of the above-stated expression (1) is set to be 0.10 or more. The .value of the above-stated expression (1) is preferably 0,20 or more, and more preferably 0,30 or more to further weaken the texture and' to randomize the crystal orientation of the austenite,
On the other hand, when the value of the above-stated expression (1) is over 1,00, the recrystaliization. of austenite excessively proceeds, the structure becomes coarse, and therefore, the value of the above-stated expression (1) is preferably 1.00 or less, and more preferably 0.90 or 1 e s s .
[0082] Besides, in the hot-rolling step, the finish rolling temperature of the hot-rolling, namely, the completion temperature of the hot-roiling is set to be 880°C or more,
When the completion temperature of the hot-rolling is less than 800°C, the development of the texture of axastenite is accelerated, the crystal orientation is strongly deflected, and there is a possibility that the crystal orientation of the retained austenite after the cold-rolling and the annealing is also deflected, Accordingly, it is important to perform the rolling at high temperature as much as possible in the hot-rolling so as not to develop the texture of the retained austenite.

n upper limit of the
On the other hand, ai
of the hot-roiling is not
completion temperature
- i A hnt- when the completion particularly provided, bu.u when tn«
;, set at an excessively high-temperature temperature is set at an
j*«« innnT it is necessary to exceeding luuu c, .1. *■■
range such as
extract the slab at a very high temperature to secure
J? ^-^ i +■ i=i not preferable
the temperature, and therefore, it .is not P
in cost phase
Accordingly, the completion
■ ferably 1000°C or less
temperature is pre [0083] (First Cooling Step)
■) 1 ■ v,^ i « -Finished, the obtained After the hot-rolling i& tinioiieu,
hot-rolled steel sheet is rapidly cooled to mate it a
coil to be a hot-rolled coil, but it is necessary to
■ * i„ control the following: an elapsed time appropriately control L.U«.
■1-ir.r, qfirts; and conditions of the until the rapid cooling starts,
uBr,1)£,e these conditions affect on the rapid cooling, because r.n«s><-
anisotropy o
present erabod
rolling completion until the rap
v. m^vr. thereafter, .the rapid jcond ox more, cm. J. «U ±. *-= /
f the hot-rolled steel sheet. In the iment, the elapsed time after the hot-
set to. be.-l. 0 se
cooling is started, a
450°C or more. Limitation reas
follows.
nd the cooling is stopped at
sons for these- are as
id cooling start is
0084] After the hot-rolling,
>lling, the texture of
u .- ,niisfi steel sheet has a-strong austenite in the hot-rolled steel
. n^nrw resulting from the processing by the anisotropy rebun.±ny
. „„4.,nr,w it is necessary rolling. To reduce the anisotropy, it
to advance the recrysta
llization of-austenite during
a period after the
hot-rolling is finished until the

subsequent rapid cooling is started, Prom this point of view, the elapsed time after the hot-rolling is completed until the rapid cooling is started is set to be 1,0 second or more. To further advance the recrystallization of austenite., the elapsed time is preferably 1.5 seconds or more, and more preferably 2,0 seconds or more. An upper limit of the. time is not particularly provided, but to start the rapid cooling after a long time over 20 seconds elapses, an enough space to retain the steel sheet after the hot-rolling is required, and facilities are necessary to be drastically increased in size. Therefore, it is not preferable in cost phase, so the time is preferably 20 seconds or less, and more preferably 15 seconds or less from a point of view of cost phase, [0085] Besides, an average cooling rate of the rapid cooling after the hot-rolling until it is coiled as a coil is preferably 10°C/second or more and 60°C/second or less. When the average cooling rate is 10°C/second or. lass, .the ferrite and the. peariite. form a ....... ... .
microstructuro extending in band-shape in a rolling direction, further Mn concentrates in. the peariite to' form an Mn concentrated region in band-shape, The retained austenite obtained by the annealing step is affected by the Mn concentrated region, easy to remain in a shape extending in the. roiling direction, and it is not preferable because there is a possibility in which the anisotropy of- the bake hardenability occurs.