High-strength hot-dip galvanized steel sheet and high-strength alloyed hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with tensile strength of 980 MPa or more and manufacturing method therefor

TECHNICAL FIELD

The present invention relates to a high-tensile steel sheet having excellent formability (ductility and hole expandability) and to an alloyed hot-dip galvanized steel sheet using TRIP (Transformation Induced Plasticity) phenomenon and a manufacturing method thereof.

This application claims priority on Japanese Patent Application No. 2011-216967, filed on Sep. 30, 2011, the content of which is incorporated herein by reference.

BACKGROUND ART

High strengthening of a steel sheet which is a raw material has been in progress so as to realize both of compatibility between a weight saving of a body, components, and the like of a vehicle, and safety. Generally, when the strength of the steel sheet increases, formability (ductility and hole expandability) is damaged. Therefore, the balance of strength and formability is necessary in order to use the high-strength steel sheet for the members of the vehicles. For this requirement, hitherto, a so-called TRIP steel sheet using transformation induced plasticity of residual austenite has been suggested (for example, refer to Patent Literature 1 and Patent Literature 2). The high strength steel sheet for the vehicle requires corrosion resistance depending on components to be applied. An alloyed hot-dip galvanized steel sheet is applied to such a case. However, Si is added to the TRIP steel to improve the ductility. When the Si concentrated on the surface of the steel sheet oxidizes, the TRIP steel has a problem in that galvanizing faults easily occur at the time of hot-dip galvanizing.

A manufacturing method of a high-strength alloyed hot-dip galvanized steel sheet has been disclosed in Patent Literatures 3 and 4 which can achieve wettability improvement of plating and reduction of alloying temperature by performing Ni pre-plating on an Si-added high strength steel sheet and working a surface layer to activate. In this method, a hot-dip galvanized steel sheet and an alloyed hot-dip galvanized steel sheet can be manufactured by re-heating and plating a cold-rolled steel sheet, in which a material as an original sheet is previously integrated, produced by a cold rolling-annealing process.

In addition, a technique has been proposed in Patent Literature 5 which produces a high elongation-type alloyed hot-dip galvanized steel sheet by utilizing a Ni pre-plating technique. This method relates to manufacturing a high-strength steel sheet having excellent corrosion resistance by making steel consisting of ferrite and martensite by controlling steel components, annealing conditions, alloying hot-dip galvanizing conditions or the like and then performing hot-dip galvanizing.

However, in the hot-dip galvanized steel sheet and the alloyed hot-dip galvanized steel sheet, the martensite, which is once generated, is softened when the steel sheet is re-heated in the galvanizing process, and thus a desired strength cannot be obtained. In this way, it is difficult to achieve both of high strengthening and formability, and a hot-dip galvanized steel sheet and an alloyed hot-dip galvanized steel sheet having good corrosion resistance with high tensile strength of 980 MPa or more and excellent formability have been desired.

PRIOR ART LITERATURE(S)
Patent Literature(s)

[Patent Literature 1] JP 61-217529A

[Patent Literature 2] JP 5-59429A

[Patent Literature 3] JP 2526320B

[Patent Literature 4] JP 2526322B

[Patent Literature 5] JP 2006-283071A

SUMMARY OF THE INVENTION
Problem(s) to Be Solved by the Invention

The present invention is to solve the above-described problems and to provide a high-strength hot-dip galvanized steel sheet and an alloyed hot-dip galvanized steel sheet having excellent ductility and hole expandability by combining a rolling process with heat treatment in a hot-dip galvanizing line.

Means for Solving the Problem(s)

The present inventors examined in detail on a structure control due to heat treatment and effects of working and heat treatment with respect to various steels in which contents of C, Si, and Mn are varied. As a result, the present inventors have found that a steel sheet having unique structures can be obtained by working and heat treatment using heat treatment in continuous annealing, rolling working, and heat treatment during subsequent plating treatment. In addition, the inventors have found that the steel sheet can have ultimate tensile strength of 980 MPa or more which has been a problem until now, excellent formability (ductility and hole expandability), and plating property.

The gist of the present invention are as follows.

[1] A high-strength hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with an ultimate tensile strength of 980 MPa or more, the hot-dip galvanized steel sheet comprising a hot-dip galvanized layer formed on a surface of a base steel sheet,

wherein the base steel sheet contains: by mass %,

C: 0.05% to 0.4%;

Si: 0.01% to 3.0%;

Mn: 0.1% to 3.0%;

Al: 0.01 to 2.0%; in which Si+Al>0.5%

P: limited to 0.04% or less;

S: limited to 0.05% or less;

N: limited to 0.01% or less; and

a balance including Fe and inevitable impurities,

a microstructure of the base steel sheet contains 40% or more by total volume fraction of martensite and bainite, 8% or more by volume fraction of residual austenite, and a balance of the microstructure being ferrite or ferrite and 10% or less by volume fraction of pearlite,

the martensite contains 10% or more by total volume fraction of two or more kinds of three kinds of martensites (1), (2), and (3) below, and

the hot-dip galvanized layer contains less than 7 mass % of Fe,

the martensite (1): C concentration (when there is a cementite precipitation, also including C in cementite); CM1 is less than 0.8 mass %, and nano-indentation test hardness Hit1 satisfies Expression 1.

Hit1/{−982.1×(CM1)²+1676×CM1+189}≤0.50 Expression 1

the martensite (2): C concentration (when there is a cementite precipitation, also including C in cementite); CM2 is 0.8 mass % or more, and nano-indentation test hardness Hit2 satisfies Expression 2.

Hit2/{−982.1×(CM2)²+1676×CM2+189}≤0.50 Expression 2

the martensite (3): C concentration (when there is a cementite precipitation, also including C in cementite); CM3 is 0.8 mass % or more, and nano-indentation test hardness Hit3 satisfies Expression 3.

0.5<Hit3/{−982.1×(CM3)²+1676×CM3+189}≤0.80 Expression 3

[2] The high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [1], wherein the base steel sheet further contains one or two or more of: by mass %,

Cr: 0.05 to 1.0%;

Mo: 0.05 to 1.0%;

Ni: 0.05 to 1.0%; and

Cu: 0.05 to 1.0%.

[3] The high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [1], wherein the base steel sheet further contains one or two or more of: by mass %,

Nb: 0.005 to 0.3%;

Ti: 0.005 to 0.3%; and

V: 0.01 to 0.5%.

[4] The high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [1], wherein the base steel sheet further contains B: 0.0001 to 0.1% by mass %.

[5] The high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [1], wherein the base steel sheet further contains one or two or more of: by mass %,

Ca: 0.0005 to 0.01%;

Mg: 0.0005 to 0.01%; and

REM: 0.0005 to 0.01%.

[6] A high-strength alloyed hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with an ultimate tensile strength of 980 MPa or more, the alloyed hot-dip galvanized steel sheet comprising an alloyed hot-dip galvanized layer formed on a surface of a base steel sheet, wherein the base steel sheet contains: by mass %,

C: 0.05% to 0.4%;

Si: 0.01% to 3.0%;

Mn: 0.1% to 3.0%;

Al: 0.01 to 2.0%; in which Si+Al>0.5%

P: limited to 0.04% or less;

S: limited to 0.05% or less;

N: limited to 0.01% or less; and

a balance including Fe and inevitable impurities,

the martensite contains 10% or more by total volume fraction of two or more kinds of three kinds of martensites (1), (2), and (3) below, and

the alloyed hot-dip galvanized layer contains 7 to 15 mass % of Fe,

the martensite (1): C concentration (when there is a cementite precipitation, also including C in cementite); CM1 is less than 0.8 mass %, and nano-indentation test hardness

Hit1 satisfies Expression 1.

Hit1/{−982.1×(CM1)²+1676×CM1+189}≤0.50 Expression 1

the martensite (3): C concentration (when there is a cementite precipitation, also including C in cementite); CM3 is 0.8 mass % or more, and nano-indentation test hardness Hit3 satisfies Expression 3.

0.5<Hit3/{−982.1×(CM3)²+1676×CM3+189}≤0.80 Expression 3

[7] The high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [6], wherein the base steel sheet further contains one or two or more of: by mass %,

Cr: 0.05 to 1.0%;

Mo: 0.05 to 1.0%;

Ni: 0.05 to 1.0%; and

Cu: 0.05 to 1.0%.

[8] The high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [6], wherein the base steel sheet further contains one or two or more of: by mass %,

Nb: 0.005 to 0.3%;

Ti: 0.005 to 0.3%; and

V: 0.01 to 0.5%.

[9] The high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [6], wherein the base steel sheet further contains B: 0.0001 to 0.1% by mass %.

[10] The high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [6], wherein the base steel sheet further contains one or two or more of: by mass %,

Ca: 0.0005 to 0.01%;

Mg: 0.0005 to 0.01%; and

REM: 0.0005 to 0.01%.

[11] A manufacturing method of a high-strength hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with an ultimate tensile strength of 980 MPa or more, the manufacturing method comprising:

with respect to a steel billet containing: by mass %,

C: 0.05% to 0.4%;

Si: 0.01% to 3.0%;

Mn: 0.1% to 3.0%;

Al: 0.01 to 2.0%; in which Si+Al>0.5%

P: limited to 0.04% or less;

S: limited to 0.05% or less;

N: limited to 0.01% or less; and

a balance including Fe and inevitable impurities,

heating to 1200° C. or higher and performing hot rolling at an Ar3 transformation temperature or higher;

performing cold rolling on a base steel sheet after the hot rolling at a reduction ratio of 40 to 70%;

annealing the base steel sheet after the cold rolling at 730 to 900° C.;

cooling the base steel sheet after the annealing to a temperature of 650 to 750° C. at an average cooling rate of 0.1 to 200° C./second, and cooling the base steel sheet to 450° C. or lower from the temperature of 650 to 750° C. at an average cooling rate of 20° C./second or faster;

holding the base steel sheet cooled to the 450° C. or lower in a range of 350 to 450° C. for 120 seconds or longer;

cooling the base steel sheet held in the range of 350 to 450° C. to 70° C. or lower at an average cooling rate of 5° C./second or faster;

rolling the base steel sheet cooled to the room temperature at an elongation percentage of 0.2 to 2%;

heating the rolled base steel sheet to “temperature of hot-dip galvanizing bath −40”° C. to “temperature of hot-dip galvanizing bath +50”° C. at an average temperature rising rate of 10° C./second or faster;

cooling the hot-dip galvanized steel sheet, which is hot-dip galvanized, to 70° C. or lower at an average cooling rate of 5° C./second or faster.

[12] The manufacturing method of the high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [11], wherein a hot-dip galvanizing bath flows at a flow rate of 10 m/min or more and 50 m/min or less at the time of the hot-dip galvanizing.

[13] The manufacturing method of the high-strength hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [11], wherein before being heated to the “temperature of hot-dip galvanizing bath −40”° C. to “temperature of hot-dip galvanizing bath +50”° C., the base steel sheet is subjected to pickling, and then a surface of the base steel sheet is polished and removed to a depth of 0.1 μm or more and is pre-plated with 0.2 to 2 g/m²of Ni.

[14] A manufacturing method of a high-strength alloyed hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with an ultimate tensile strength of 980 MPa or more, the manufacturing method comprising:

with respect to a steel billet containing: by mass %,

C: 0.05% to 0.4%;

Si: 0.01% to 3.0%;

Mn: 0.1% to 3.0%;

Al: 0.01 to 2.0%; in which Si+Al>0.5%

P: limited to 0.04% or less;

S: limited to 0.05% or less;

N: limited to 0.01% or less; and

a balance including Fe and inevitable impurities,

heating to 1200° C. or higher and performing hot rolling at an Ar3 transformation temperature or higher;

performing cold rolling on a base steel sheet after the hot rolling at a reduction ratio of 40 to 70%;

annealing the base steel sheet after the cold rolling at 730 to 900° C.;

holding the base steel sheet cooled to the 450° C. or lower in a range of 350 to 450° C. for 120 seconds or longer;

cooling the base steel sheet held in the range of 350 to 450° C. to 70° C. or lower at an average cooling rate of 5° C./second or faster;

rolling the base steel sheet cooled to the room temperature at an elongation percentage of 0.2 to 2%;

dipping and hot-dip galvanizing the base steel sheet heated to the “temperature of hot-dip galvanizing bath −40”° C. to “temperature of hot-dip galvanizing bath +50”° C. into a hot-dip galvanizing bath and performing alloying-heating treatment at “temperature of hot-dip galvanizing bath −40”° C. or higher and 560° C. or lower within 40 seconds; and

cooling the alloyed hot-dip galvanized steel sheet, which is subjected to the alloying-heating treatment, to 70° C. or lower at an average cooling rate of 5° C./second or faster.

[15] The manufacturing method of the high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [14], wherein a hot-dip galvanizing bath flows at a flow rate of 10 m/min or more and 50 m/min or less at the time of the hot-dip galvanizing.

[16] The manufacturing method of the high-strength alloyed hot-dip galvanized steel sheet having the excellent plating adhesion, formability, and hole expandability with the ultimate tensile strength of 980 MPa or more according to [14], wherein before being heated to the “temperature of hot-dip galvanizing bath −40”° C. to “temperature of hot-dip galvanizing bath +50”° C., the base steel sheet is subjected to pickling, and then a surface of the base steel sheet is polished and removed to a depth of 0.1 μm or more and is pre-plated with 0.2 to 2 g/m²of Ni.

Effects(S) of the Invention

According to the present invention, it is possible to obtain a hot-dip galvanized steel sheet and an alloyed hot-dip galvanized steel sheet having excellent formability with high strength and to remarkably contribute to the industry.

MODE(S) FOR CARRYING OUT THE INVENTION

The present invention will be described below in detail.

Microstructure of Base Steel Sheet

First, microstructures of a base steel sheet according to the present invention will be described. The microstructures of the base steel sheet according to the present invention include bainite, martensite, and a residual austenite. Further, the martensite includes two or more kinds of three kinds of martensites (1), (2), and (3) defined below.

Martensite (1): C concentration (when there is a cementite precipitation, also including C in cementite). CM1 is less than 0.8 mass % and nano-indentation test hardness Hit1 satisfies Expression 1.

Hit1/{−982.1×(CM1)²+1676×CM1+189}≤0.50 Expression 1

Martensite (2): C concentration (when there is a cementite precipitation, also including C in cementite); CM2 is 0.8 mass % or more, and nano-indentation test hardness Hit2 satisfies Expression 2.

Hit2/{−982.1×(CM2)²+1676×CM2+189}≤0.50 Expression 2

Martensite (3): C concentration (when there is a cementite precipitation, also including C in cementite); CM3 is 0.8 mass % or more, and nano-indentation test hardness Hit3 satisfies Expression 3.

0.5<Hit3/{−982.1×(CM3)²+1676×CM3+189}≤0.80 Expression 3

Although the detailed reason is unclear, when two or more kinds of martensites among these three kinds of martensites (1) to (3) are contained 10% or more by total volume fraction, both of strength and hole expandability are achieved. The martensite, which has the hardest structure of structures contained in the base steel sheet of the present invention, is essential to ensure ultimate tensile strength of 980 MPa or more. On the other hand, in a hole expanding test and a bending test, since the martensite is a starting point of void formation, it is known that the martensite deteriorates hole expandability. Therefore, in order to ensure the hole expandability, deformation is prevented from concentrating on a specific martensite grain by mixing two or more kinds of martensites among these three different kinds of martensites. As a result, a technique that does not deteriorate the hole expandability while contributing to higher strengthening has been found. This effect can be obtained when the total volume fraction of the two or more martensites among the three kinds of martensites is set to be 10% or more. From this reason, the lower limit of the volume fraction of the three kinds of martensites (1) to (3) was set to be 10%. Preferably, the lower limit is 15% or more.

The martensite (1) is a tempered martensite, which is low in C concentration and is not so hard. The C concentration CM1 of the martensite (1) is less than 0.8 mass %. When there is a cementite precipitation in the martensite (1), C in cementite which is precipitated in the martensite (1) is also contained. The cementite in the martensite mentioned herein may be either of a precipitation in or between martensite laths, so that the effect of the present invention is exhibited. This structure originates from a fresh martensite to be formed at the time of cooling to 70° C. or lower after being held at 350 to 450° C. in an annealing process of the manufacturing method of the present invention. The martensite (1) is a tempered martensite in which the fresh martensite formed at the time of cooling to 70° C. or lower after being held at 350 to 450° C. is tempered during dipping treatment into a hot-dip galvanizing bath or alloying treatment.

The C concentration CM1 of the martensite (1) is less than 0.8 mass %. This also includes a case in which C concentration was reduced to less than 0.8 mass % when C in the fresh martensite was diffused into austenite during the dipping treatment into the hot-dip galvanizing bath or the alloying treatment and thus the fresh martensite was tempered even though C concentration of the fresh martensite was 0.8 mass % or more, in addition to a case in which C concentration of the fresh martensite formed at the time of cooling to 70° C. or lower after being held at 350 to 450° C. was less than 0.8 mass %.

Since the martensite (1) has the C concentration CM1 as low as less than 0.8 mass % and is tempered, it is the softest among three kinds of martensites (1) to (3). Nano-indentation test hardness Hit1 of the martensite (1) satisfies Expression 1. When the volume fraction of the martensite (1) is 60% or more of a microstructure of the base steel sheet, a volume fraction of ferrite and residual austenite becomes too low, and the ductility deteriorates, so that the upper limit is preferably 60%.

The martensite (2) has a high C concentration, but is martensite which is softened by tempering. C concentration CM2 of the martensite (2) is 0.8 mass % or more. When there is a cementite precipitation in the martensite (2), C in the cementite which is precipitated in the martensite (2) is also contained. Similarly, the cementite in the martensite may be either of precipitation in or between martensite laths. The martensite (2) originates from strain-induced transformation martensite which is induced by the rolling work and into which a part of a residual austenite is transformed. While the base steel sheet is held at a temperature range of 350 to 450° C. after annealing, bainite transformation of the microstructure in the base steel sheet proceeds and thus C is diffused into a non-transformed austenite. A residual austenite, in which C is concentrated, is formed in the base steel sheet which is cooled to 70° C. or lower. By strain-induced transformation of the residual austenite, in which C is concentrated, by rolling, martensite which is the origin of the martensite (2) can be obtained. The Martensite (2) is a tempered martensite in which the strain-induced transformation martensite generated by the rolling is tempered during the dipping treatment into the hot-dip galvanizing bath or the alloying treatment.

Like the martensite (1), the martensite (2) is a tempered martensite, but is harder than the martensite (1) since the C concentration is high. Nano-indentation test hardness Hit2 of the martensite (2) satisfies Expression 2. When a volume fraction of the martensite (2) is 40% or more of the microstructure of the base steel sheet, the ductility deteriorates, so that the upper limit is preferably set to be 60%.

The martensite (3) is martensite which is high in C concentration, is not tempered, or is low in a quantity of tempering. C concentration CM3 of the martensite (3) is 0.8 mass % or more. When there is a cementite precipitation in the martensite (3), C in the cementite which is precipitated in the martensite (3) is also contained. Similarly, the cementite in the martensite may be either of precipitation in or between martensite laths. The martensite (3) is a fresh martensite which is formed by transforming at the time of cooling to 70° C. or lower after the hot-dip galvanizing or the alloying-heating treatment. In addition, the martensite (3) is also martensite remaining in a state of fresh martensite without being tempered substantially during the alloying-heating treatment or the alloying treatment (not in a state of martensites (1) and (2)).

The martensite (3) is the hardest structure of three kinds of martensites (1) to (3). Nano-indentation test hardness Hit3 of the martensite (3) satisfies Expression 3. For this reason, the martensite (3) contributes to high strengthening, but deteriorates the hole expandability. Accordingly, in order to achieve both of the strength and formability, the upper limit of the martensite (3) is 10% in the volume fraction of the microstructure of the base steel sheet. However, when the volume fraction of the martensite (3) becomes too small, it is difficult to ensure strength which is ultimate tensile strength of 980 MPa or more, so that the lower limit is preferably set to be 3% or more.

Bainite is effective for ensuring the strength. When a high strength steel sheet having tensile strength exceeding 980 MPa contains martensite and bainite of 40% or more by total volume fraction, it is possible to ensure the strength and to obtain the high hole expandability. When the total volume fraction is less than 40%, the tensile strength is less than 980 MPa. Accordingly, the lower limit was set to be 40%.

The residual austenite is a structure to raise ductility, particularly uniformly elongation by transformation-induced plasticity. In order to obtain particularly good elongation, it is necessary to contain the residual austenite of 8% or more by volume fraction. Furthermore, due to transform into martensite by working, the residual austenite also contributes to obtaining of the high strengthening.

In the microstructure of the base steel sheet of the present invention, ferrite is not essential. However, since the ferrite causes the improvement of ductility, it may be contained. At the time of annealing, it is possible to control the volume fraction of ferrite by performing two-phase annealing. Furthermore, it is possible to control the volume fraction of ferrite by cooling after annealing. However, when the fraction of ferrite increases, the strength decreases. Although the high strengthening can be obtained by precipitation strengthening and solid solution strengthening, the volume fraction of ferrite is preferably 40% or less.

Pearlite may be contained as long as the volume fraction is 10% or less. When the volume fraction of pearlite exceeds 10%, the strength and ductility are reduced. Therefore, the upper limit was set to be 10%.

Furthermore, each phase of the microstructures such as martensite, bainite, austenite, pearlite, and ferrite can be identified and their locations and volume fraction can be observed and quantitatively measured using an optical microscope having a magnification of 1000 times and a scanning and transmission electron microscope having a magnification of 1000 times to 100000 times after a cross section of the steel sheet in a rolling direction or a cross section in the right angle direction of the rolling direction is etched using a Nital reagent and the reagent as disclosed in JP 59-219473A. The area fraction of each structure can be obtained by observing 20 or more fields and applying the point-count method or image analysis. Then, the obtained area fraction is defined as the volume fraction of each structure.

The classification method of three kinds of martensites (1) to (3) is hardness and C concentration. The hardness may be obtained by measuring nano-indentation testing hardness for three or more points in martensite grains and calculating an average hardness Hit. In the base steel sheet according to the present invention, since a large amount of additive elements are contained, a crystal grain diameter is small. Moreover, in the base steel sheet according to the present invention, there is a case in which an indentation size is greater than the grain diameter of martensite grains in measurement using a Vickers hardness test. Therefore, hardness measurement of a micro-region was performed by a nano-indenter. Samples cut out parallel to the rolling direction are embedded and then are subjected to mirror-polishing and electrolytic polishing. Then, the hardness measurement was performed on the polished samples. As test conditions, an indentation depth was measured under the condition of 50 nm using a Berkovich-type indenter. Furthermore, this test method is small in an indentation depth, and is sensitive to relation between the grain diameter of martensite and the indentation size, or surface irregularities. Therefore, as a preliminary test, the electrolytic polishing in various conditions and the hardness measurement in the condition with varying an indentation depth are performed, and conditions for obtaining a most reproducible good value was set as test conditions.

The C concentration of martensite grains may be measured by any measuring method which guarantees precision under the condition of obtaining an accurate decomposition concentration. For example, the C concentration of martensite grains can be obtained by carefully measuring the C concentration at a pitch of 0.5 μm or below using EPMA attached to FE-SEM. Therefore, martensites (1) to (3) are classified according to the hardness and the C concentration.

Furthermore, in order to distinguish these martensites (1) to (3), Expressions 1 to 3 use a relational expression between amounts of C CM1 to CM3 and nano-indentation test hardness Hit in each martensite. Denominators of the left sides of Expressions 1, 2, and 3, which are input values of C concentration represent hardness of carbides-free martensite (fresh martensite) of the C concentration. The hardness of the martensite contained in the base steel sheet of the present invention becomes lower than the hardness of a fresh martensite by precipitation of cementite at grains and tempering. Therefore, the classification is performed by taking a ratio between hardness of a fresh martensite of denominator and hardness of martensite of a steel sheet.

Chemical Composition of Base Steel Sheet

Next, a description of reasons for restricting the amounts of the compositions of the base steel sheet in the present invention will be described. Moreover, % in the composition represents % by mass.

C: C is added as an element which increases strength of steel and stabilizes a residual austenite to improve ductility. When the content is less than 0.05%, it is difficult to ensure tensile strength of 980 MPa or more. Ductility, weldability, and toughness are remarkably deteriorated by excessive addition exceeding 0.40%. Therefore, the content of C was set to be 0.05 to 0.4%. A more preferred range is 0.13% to 0.3%,

Si: Si is an element useful for increasing the strength of the steel sheet by solid-solution strengthening. In addition, Si is an essential element which has an effect for promoting concentration of C in austenite during bainite transformation and generates a residual austenite while annealing, in order to suppress the formation of cementite. These effects are not exhibited when the content is less than 0.01% and scale exfoliation and chemical convertibility generated in hot rolling are remarkably deteriorated by excessive addition exceeding 3.0%. Thus, the content of Si was set to be 0.01 to 3.0%.

Mn: Mn is an element effective for improving hardenability. An effect of increasing the hardenability is not sufficiently exhibited when the content is less than 0.1% and the toughness is deteriorated by excessive addition exceeding 3.0%. Accordingly, the content of Mn was set to be 0.1 to 3.0%.

Al: Al is an element having a function of deoxidizer. In addition, Al is a ferrite stabilizing element like Si and may be also used as an alternative of Si. Such an effect is not exhibited when the content is less than 0.01% and the toughness is deteriorated by excessive addition exceeding 2.0%. Therefore, the content of Al was set to be 0.01 to 2.0%.

Al+Si: Al and Si are elements having the same functions of ferrite stabilization and cementite suppression. Accordingly, a total additive amount of Al and Si is important. When the total additive amount is 0.5% or less, the function of stabilizing the ferrite and suppressing the cementite becomes weak. Therefore, the content was added with an amount larger than 0.5%.

P: P is an impurity element which segregates at grain boundaries to make grain boundary strength lower, thereby deteriorating the toughness. Thus, the content is preferably reduced. The upper limit of the content of P was limited to 0.04% in consideration of a current refining technology and manufacturing costs. The lower limit value of P is not particularly determined, but when the lower limit value is less than 0.0001%, it is disadvantageous economically, so this value is preferably set to the lower limit value.

S: S is an impurity element which deteriorates hot workability and toughness, and the content is preferably reduced. Accordingly, the upper limit was limited to 0.05%. The lower limit value of S is not particularly determined, but when the lower limit value is less than 0.0001%, it is disadvantageous economically, so this value is preferably set to the lower limit value.

N: N forms coarse nitrides to deteriorate bendability and hole expandability. Therefore, it is necessary to suppress the additive amount. The reason is because the tendency becomes remarkable when the content of N exceeds 0.01%. Thus, the content of N was in a range of 0.01% or less. In addition, this causes blowholes to occur at the time of welding, so the less the better. The effect of the present invention is exhibited without particularly determining the lower limit, but when the content of N is less than 0.0005%, the manufacturing cost dramatically increases, so this value is a substantial lower limit.

Further, one or two or more elements of Cr, Mo, Ni, and Cu may be added. These elements are elements effective for improving ductility and toughness. However, when the content of Cr, Mo, Ni, and Cu exceeds 1.0%, the toughness can be impaired due to an increase in strength. Accordingly, the upper limit of these elements was set to be 1.0%. Further, in order to improve the ductility and toughness, the necessary content of Cr is 0.05% or more, the necessary content of Mo is 0.05% or more, the necessary content of Ni is 0.05% or more, and the necessary content of Cu is 0.05% or more, so these values are set to the lower limit value, respectively.

Further, one or two or more elements of Ti, Nb, and V may be added. These elements are elements which form fine carbonitrides and are effective for suppressing coarsening of crystal grains, ensuring the strength, and improving the toughness. In order to ensure the strength and to improve the toughness, it is necessary to add 0.005% or more Ti and Nb and 0.01% or more V. However, when these elements are excessively added, a precipitate becomes coarse and the toughness may be deteriorated. Accordingly, the additive amount of Nb and Ti is preferably set to be 0.3% or less, and the additive amount of V is preferably set to be 0.5 or less.

B: B is an element which segregates at grain boundaries to suppress grain boundary segregation of P and S. In addition, this element is also effective for improving the hardenability. However, when the content of B exceeds 0.1%, a coarse precipitate occurs at the grain boundaries to impair the hot workability and the toughness. Accordingly, the content of B is set to be 0.1% or lower. Further, in order to enhance the ductility, toughness, and hot workability and to improve the hardenability by the strengthening of the grain boundaries, the addition of B is preferably 0.0001% or more.

Further, one or two or more elements of Ca, Mg, and REM may be added. These elements are elements effective for controlling sulfide forms to suppress the deterioration of the hot workability and toughness due to S. REM indicates a rare earth metal. However, even when these elements are excessively added, since the effect is saturated, it is preferable that 0.01% or less Ca, 0.01% or less Mg, and 0.01% or less REM be added, respectively. In order to improve the toughness, 0.0005% or more Ca, 0.0005% or more Mg, and 0.0005% or more REM are preferably added, respectively. Further, in the present invention, REM is generally added in a mischmetal, which in addition to La and Ce may also contain other lanthanoid series elements in combination. The effects of the invention are exhibited even when the lanthanoid series elements other than La and Ce are contained as inevitable impurities. However, the effects of the present invention are exhibited even when metals such as La and Ce are added.

Chemical Composition of Hot-Dip Galvanized Layer and Alloyed Hot-Dip Galvanized Layer

In the present invention, a hot-dip galvanized layer formed on the surface of the base steel sheet contains less than 7 mass % Fe, the balance being Zn and inevitable impurities. In addition, an alloyed hot-dip galvanized layer contains 7 to 15 mass % Fe, and the balance being Zn and inevitable impurities. The hot-dip galvanized layer and the alloyed hot-dip galvanized layer may further contain Al of 0.01 to 0.5 mass % and more preferably, may contain Al of 0.05 to 0.3 mass %. Further, the galvanizing bath may contain Fe, Mg, Mn, Si, Cr and the like in addition to pure zinc and Al.

In a case where spot weldability and a coating property are desired, it is possible to improve these properties by forming the alloyed hot-dip galvanized layer containing 7 to 15 mass % Fe on the surface of the base steel sheet. Specifically, when the base steel sheet is subjected to the alloying treatment while being dipped in the galvanizing bath, Fe is incorporated into the galvanized layer, and thus the high-strength alloyed hot-dip galvanized steel sheet having an excellent coating property and spot weldability can be obtained. When the content of Fe after the alloying treatment is less than 7 mass %, the spot weldability becomes insufficient. On the other hand, when the content of Fe exceeds 15 mass %, the adhesion of the galvanized layer itself is impaired, and the galvanized layer is broken and dropped out in machining, thereby causing scratches when forming by adhering to a mold. Accordingly, the content of Fe contained in the galvanized layer during the alloying treatment is within a range of 7 to 15 mass %.

Further, in a case where the alloying treatment is not performed, even when the content of Fe contained in the galvanized layer is less than 7 mass %, the corrosion resistance, the formability, and hole expandability which are effects obtained by the alloying are good except for the spot welding.

Further, the galvanized layer may contain Al, Mg, Mn, Si, Cr, Ni, Cu or the like in addition to Fe.

In order to measure the content of Fe and Al contained in the galvanized layer, a method of dissolving the galvanized layer with an acid and chemically analyzing the dissolved solution may be used. For example, with respect to the alloyed hot-dip galvanized steel sheet cut into 30 mm×40 mm, only the galvanized layer is dissolved while suppressing elution of the base steel sheet with an inhibitor-added 5% HCl aqueous solution. Then, the content of Fe and Al is quantified using signal intensities obtained by ICP emission analysis of the dissolved solution and a calibration curve prepared from concentration-known solutions. Further, in consideration of measured variation of samples, an average value is employed obtained by measuring at least three samples which are cut out from the same alloyed hot-dip galvanized steel sheet.

The coated amount of the plating is not particularly limited, but is preferably 5 g/m²or more in the coated amount on a single surface of the base steel sheet from the viewpoint of corrosion resistance. In addition, the coated amount on the single surface is preferably no greater than 100 g/m²from the viewpoint of ensuring the plating adhesion.

Manufacturing Method of Steel Sheet

Next, a manufacturing method will be described.

In the present invention, the steel consisting of the above compositions is casted by melting in a conventional manner. The obtained steel billet (cast slab) is subjected to hot rolling. The cast slab is directly cooled or once cooled and then is heated to 1200° C. or higher, and the hot rolling is finished at an Ar3 transformation temperature or higher.

The base steel sheet (hot-rolled steel sheet) subjected to the hot rolling may be coiled at a temperature region of 700° C. or lower. When the coiling temperature exceeds 700° C., the structure of the hot-rolled steel sheet becomes a coarse ferrite or pearlite structure. As a result, a structure of a final steel sheet becomes a non-uniform structure, and thus it is difficult to obtain good hole expandability. Therefore, the upper limit of the coiling temperature is set to be 700° C. More preferably, the upper limit is 650° C. or lower and, most preferably, is 550° C. at which bainite-single phase is formed. The lower limit of the coiling temperature is not particularly defined. However, when the lower limit is lower than 300° C., the strength of the hot-rolled sheet increases and causes interference of a cold rolling in some cases. Therefore, the lower limit is preferably 300° C. or higher.

Then, the base steel sheet (hot-rolled steel sheet) is subjected to pickling treatment as necessary and then is subjected to the cold rolling at a reduction ratio of 40 to 70%. In order to refine a microstructure after annealing, the cold rolling is performed at the reduction ratio of 40% or more. On the other hand, when the reduction ratio of the cold rolling exceeds 70%, a load is increased by work hardening to cause a loss of productivity. Accordingly, the reduction ratio of the cold rolling is set to be 40 to 70%.

After the cold rolling, the base steel sheet is annealed at 730 to 900° C. In order to control the microstructure of the base steel sheet, a heating temperature of the annealing and cooling conditions are very important in the present invention. The annealing after the cold rolling is performed at the range of 730 to 900° C. to obtain austenite in which C is sufficiently concentrated. When the annealing temperature is lower than 730° C., carbides are melted and remain, and thus a required amount of austenite cannot be obtained. When the annealing temperature exceeds 900° C., it is uneconomical. Further, in a case where the annealing temperature exceeds 900° C., recrystallization proceeds and a grain diameter becomes larger, thereby deteriorating toughness and ductility. Therefore, the annealing temperature is set to be 730 to 900° C.

After being subjected to the annealing, the base steel sheet is cooled to a temperature of 600 to 750° C. at an average cooling rate of 0.1 to 200° C./second. Thereafter, the base steel sheet is cooled to a temperature of 450° C. or lower from a temperature of 600 to 750° C. at an average cooling rate of 20° C./second or faster. The purpose of cooling the base steel sheet to the temperature of 600 to 750° C. at the average cooling rate of 0.1 to 200° C./second from the temperature region of the annealing is to suppress formation of pearlite which occurs during the cooling process. When the cooling rate is slower than 0.1° C./second, it is difficult to avoid pearlite transformation and a part or all of the austenite is transformed into the pearlite, so that it is difficult to obtain the high strength of 980 MP or more. Therefore, the average cooling rate from the temperature region of the annealing to the temperature of 600 to 750° C. is set to be 0.1° C./second or faster. On the other hand, the cooling at the cooling rate of 200° C./second or faster causes not only saturation of the effect but also excessive facility investment, so that economic efficiency becomes poor. Therefore, the upper limit of the cooling rate is set to be 200° C./second.

When the cooling rate from the temperature of 600 to 750° C. to the temperature of 450° C. or lower is slow, bainite transformation proceeds and a large amount of carbides are formed in the bainite structure, so that the austenite is decomposed and thus the ductility becomes weak. In addition, since three kinds of martensites (1) to (3) may not be obtained, the balance of the strength and the hole expandability is low. For this reason, it is necessary to cool at the average cooling rate of 20° C./second or faster. Furthermore, the upper limit is not limited, but when the cooling rate is excessively raised, manufacturing cost increases, and thus the upper limit is preferably 200° C./second or slower.

Next, the base steel sheet is held at the range of 350 to 450° C. for 120 seconds or longer. Further, as described above, when the cooling is performed from the temperature of 600 to 750° C. to the temperature of 450° C. or lower, a cooling stop temperature is set to be 350° C. or higher, and then the base steel sheet may be held at the range of 350 to 450° C. for 120 seconds or longer. Alternatively, when the cooling is performed from the temperature of 600 to 750° C. to the temperature of 450° C. or lower, the cooling stop temperature is set to be lower than 350° C., and then the base steel sheet is re-heated and may be held at the range of 350 to 450° C. for 120 seconds or longer. The reason is because the holding at the temperature region of 350 to 450° C. is to control the stability of austenite. When the holding time is short, the stabilization of austenite is not achieved and a residual austenite 8% or more by volume fraction cannot be obtained. As a result, the balance of the strength and the ductility will be failed. On the other hand, the holding for a long time causes excessive stabilization of the residual austenite, and thus martensite is not formed in the cooling process to room temperature to be subsequently performed and three kinds of martensites (1) to (3) are not obtained. Accordingly, the holding time is preferably set to be 1000 seconds or shorter.

Subsequently, the base steel sheet is cooled once to 70° C. or lower at the average cooling rate of 5° C./second or faster. A cooling attainment temperature may be room temperature. However, since the attainment temperature varies depending on the season, the attainment temperature may be 70° C. or lower in terms of securing materials and may vary depending on the season. In addition, the average cooling rate up to 70° C. needs to be set to be 5° C./second or faster. The upper limit is not particularly determined, but the cooling at the cooling rate over 200° C./second causes not only saturation of the effect but also large facility investment, so that economic efficiency becomes poor. Therefore, the upper limit of the cooling rate is preferably set to be 200° C./second or slower. For this reason, a part of austenite is transformed into martensite. By performing such heat treatment, a composite structure of bainite, a residual austenite, and martensite can be obtained. However, ferrite may be partially contained. Further, the martensite obtained by the cooling up to the room temperature is a fresh martensite which is the origin of the martensite (1).

Next, the base steel sheet cooled to 70° C. or lower is rolled. A part of residual austenite in the base steel sheet is transformed by the rolling working, and thus a strain-induced transformation martensite is generated. Thus, the strain-induced transformation martensite obtained by the rolling is the origin of the martensite (2). The austenite remaining at the room temperature contains C of 0.8 mass % or more, and the strain-induced transformation martensite formed by the transformation of the above austenite becomes hard. In order to obtain a sufficient strain-induced transformation martensite, an elongation percentage (reduction ratio) of the rolling is set to be 0.2 to 2%. The effect is not sufficient when the elongation percentage is less than 0.2%, and the yield ratio is significantly increases and the elongation deteriorates when the elongation percentage exceeds 2%.

Next, after being subjected to Ni pre-plating as necessary, the base steel sheet is subjected to galvanizing or galvanizing and alloying-heating treatment. In the case of performing the Ni pre-plating, after the base steel sheet is subjected to the pickling, the surface of the base steel sheet is removed by polishing to 0.1 μm or more in depth and then Ni is pre-plated on the surface of the base steel sheet with 0.2 to 2 g/m². In order to suppress of galvanizing faults and perform the alloying, the surface of the steel sheet is preferably subjected to the polishing or the Ni pre-plating. The reason is because oxides are easily formed on the surface of the base steel sheet and the galvanizing faults and the alloying treatment delay of the hot-dip galvanizing easily occur through a plurality of heating treatment processes.

After the annealing, oxides of Si, Mn and the like exist in the surface layer of the base steel sheet in some cases. When these oxides exist, even if the base steel sheet is assumed to be subjected to the Ni pre-plating, the galvanizing faults occur in some cases. For this reason, it is necessary to remove the oxides by performing the polishing. Since the effect becomes remarkable by polishing to 0.1 μm or more in depth from the surface layer of the base steel sheet, a polishing quantity is set to be 0.1 μm. The upper limit is not particularly determined. However, since the steel sheet becomes thin and product yield is reduced according to the polishing quantity, the polishing quantity is preferably small.

When a Ni-coated amount is set to be 0.2 to 2 g/m², the galvanizing faults are suppressed at the time of subsequent hot-dip galvanizing. In the Ni-coated amount of less than 0.2 g/m², since the effect of suppressing the galvanizing faults is not sufficiently obtained, the lower limit thereof is set to be 0.2 g/m². The effect of suppressing the galvanizing faults can be obtained without specifically determining the upper limit, but the upper limit exceeding 2.0 g/m²causes not only saturation of the effect but also excessive facility investment, so it is undesirable that the economic efficiency becomes poor. In addition, this requires the excessive facility investment or an operation dropped extremely in sheet-passing speed is required, so it is undesirable that the economic efficiency becomes poor.

A temperature of the base steel sheet to be dipped in the hot-dip galvanizing bath is in a range from a temperature that is lower by 40° C. compared with the temperature of the hot-dip galvanizing bath to a temperature that is higher by 50° C. compared with the temperature of the hot-dip galvanizing bath. When the temperature of the base steel sheet to be dipped is below “temperature of hot-dip galvanizing bath −40” ° C., the heat loss upon dipping into the galvanizing bath becomes large and a part of the molten zinc is solidified, thereby leading to a deterioration of the galvanized external appearance in some cases. In addition, when the temperature of the base steel sheet is above “temperature of hot-dip galvanizing bath +50”° C., operational problems associated with a temperature rise of the galvanizing bath are induced. Further, the temperature of the galvanizing bath is managed to 440 to 470° C. The decrease in temperature of the galvanizing bath leads to solidification of the hot-dip galvanizing existing in the bath and becomes a cause of the galvanizing faults or becomes a cause of appearance deterioration.

Thus, the rolled base steel sheet should be heated to the “temperature of hot-dip galvanizing bath −40”° C. to the “temperature of hot-dip galvanizing bath +50”° C. Here, the base steel sheet is heated at an average temperature rising rate of 10° C./second or higher. Here, when the heating rate is slower than 10° C./second, the surface of the base steel sheet is polished and removed while the Ni pre-plating and the induced strain is relaxed, so that alloying-promoting effect cannot be obtained. In addition, when the heating temperature is lower than the “temperature of hot-dip galvanizing bath −40”, the galvanizing faults easily occur during the hot-dip galvanizing. When the heating temperature is above the “temperature of hot-dip galvanizing bath +50”° C., the surface of the base steel sheet is polished and removed and the induced strain is relaxed, so that the alloying-promoting effect cannot be obtained.

In a hot-dip galvanizing tank, it is preferable that a jet flow of 10 m/min or more and 50 m/min or less be provided in the galvanizing bath to suppress the galvanizing faults and to promote the alloying. Scum, which is an oxide film of Zn or Al, is floated on the surface of the galvanizing bath. When the oxide film remains on the surface of the base steel sheet in large amounts, the scum adheres to the surface of the base steel sheet at the time of dipping in the galvanizing bath and the galvanizing faults easily occur. Further, the scum adhering to the steel sheet causes not only the galvanizing faults but also the alloying delay.

This property is particularly remarkable in the steel sheet containing a lot of Si and Mn. The detailed mechanism is unclear, but it is considered that the galvanizing faults and the alloying delay are facilitated by reacting between the oxide of Si or Mn, which is formed on the surface of the base steel sheet, and the scum that is similarly the oxide. The reason for setting the flow rate of the jet flow to be 10 m/min or more and 50 m/min or less is because the suppressing effect of the galvanizing faults due to the jet flow cannot be obtained at the flow rate slower than 10 m/min. The reason for setting the flow rate to be 50 m/min or less is because the suppressing effect of the galvanizing faults is saturated and a high cost due to the excessive facility investment is also avoided.

In addition, the galvanizing bath may contain Fe, Al, Mg, Mn, Si, Cr and the like in addition to pure zinc.

Then, when the base steel sheet is subjected to the hot-dip galvanizing or the hot-dip galvanizing and the alloying treatment by dipping in the hot-dip galvanizing bath, the martensite in the base steel sheet is tempered. That is, as described above, the fresh martensite which is the origin of the martensite (1) and the strain-induced transformation martensite which is the origin of the martensite (2) are contained in the base steel sheet to be dipped into the hot-dip galvanizing bath. When the base steel sheet is subjected to the dipping and the alloying heat treatment in the hot-dip galvanizing bath, the fresh martensite and the strain-induced transformation martensite formed previously in the base steel sheet are tempered. As a result, the martensite (1) and the martensite (2) are formed.

Next, the hot-dip galvanized steel sheet subjected to the hot-dip galvanizing or the alloyed hot-dip galvanized steel sheet subjected to the hot-dip galvanizing and the alloying treatment is cooled to 70° C. or lower. The cooling attainment temperature may be set to be room temperature. However, since the attainment temperature varies depending on the season, the attainment temperature may be 70° C. or lower in terms of securing materials and may vary depending on the season. In addition, the average cooling rate up to 70° C. needs to be set to be 5° C./second or faster. The upper limit is not particularly determined, but the cooling at the cooling rate over 200° C./second causes not only the saturation of the effect but also the large facility investment, so that it is economically undesirable. Therefore, the upper limit of the cooling rate is preferably set to be 200° C./second or slower. By the cooling, a part of residual austenite in the base steel sheet is transformed into martensite. The martensite generated in this way becomes a martensite (3) which has the hardest structure.

In only the heat treatment, only the decomposition of the residual austenite is caused. However, in the present invention, the steel sheet is previously subjected to the rolling and thus the residual austenite is further processed. As a result, a part of residual austenite is transformed into martensite while the cooling after the heat treatment. The martensite (3) obtained as a result becomes a reinforcing structure to achieve high strengthening. The detailed mechanism is unclear, but when the structure of the steel sheet subjected to the rolling was compared with the structure of the steel sheet, which was not subjected to temper rolling, in detail, this martensite was not observed in the steel sheet which was not subjected to the rolling. For this reason, it is assumed that a dislocation induced while the rolling contributes to the martensite transformation at the time of the heat treatment and subsequent treatment.

Furthermore, in order to obtain the martensite (3), it is necessary to heat the base steel sheet to the temperature of “temperature of hot-dip galvanizing bath −40” (° C.) or higher and 560° C. or lower. In the heating at the “temperature of hot-dip galvanizing bath −40” (° C.) or lower, the martensite cannot be obtained while the cooling to be performed subsequently. On the other hand, when the base steel sheet is heated to the temperature exceeding 560° C., carbides are remarkably precipitated and the austenite is decomposed, and thus it is difficult to obtain the amount of residual austenite which is required for elongation improvement. For this reason, in the case of performing the hot-dip galvanizing, the heating temperature of the base steel sheet is set to be 560° C. or lower. In addition, when the time required for the alloying heat treatment is long, the austenite is decomposed. Accordingly, an alloying treatment time is preferably set to be 40 seconds or shorter.

Further, in a facility such as a continuous hot-dip galvanizing facility for continuously performing the heat treatment and plating, the steel sheet is not cooled to the room temperature in a middle course and is not subjected to the rolling in a middle course. Consequently, structure controls of separately manufacturing three kinds of martensites as in the present invention and of ensuring the residual austenite could not be performed. Accordingly, it was difficult to achieve all of the strength, ductility, and hole expandability with a high level.

In addition, when the surface layer of the base steel sheet is pre-plated with Ni after being polished and removed to 0.1 μm or thicker, the alloying is promoted during the alloying-heating treatment after the galvanizing to lower the heating temperature at the time of the alloying treatment. For this reason, the cementite is not generated during the alloying-heating treatment and the deterioration of the hole expandability is avoided. The mechanism of the alloying promotion is unclear, but it is considered that the surface is activated due to the strain induced into the surface layer of the steel sheet by the polishing. Examples of methods of polishing and removing the surface layer of the base steel sheet may include brush polishing, sandpaper polishing, or mechanical polishing. The method of the Ni pre-plating may be any one of electroplating, dipping-plating, and spraying-plating, and plating weight is preferably about 0.2 to 2 g/m². When the polishing and removal amount of the surface layer of the steel sheet is 0.1 μm or less and the Ni pre-plating is not performed, or when the pre-plating weight is 0.2 g/m²or less or 2 g/m², the promotion effect of the alloying is not obtained and the alloying temperature is inevitably raised. Thus, as described below, the deterioration of the hole expandability is not prevented. In order to further obtain the promotion effect of the alloying, the polishing and removal amount of the surface layer of the steel sheet is preferably set to be 0.5 μm or more.

Further, in the case of manufacturing the alloyed hot-dip galvanized steel sheet, an effective Al concentration in the galvanizing bath is preferably controlled in the range of 0.05 to 0.500 mass % to control the properties of the galvanized layer. Here, the effective Al concentration in the galvanizing bath is a value obtained by subtracting a Fe concentration in the galvanizing bath from the Al concentration in the galvanizing bath.

When the effective Al concentration is less than 0.05 mass %, dross significantly occurs and a good appearance cannot be obtained. On the other hand, the effective Al concentration is more than 0.500 mass, the alloying is delayed and the productivity is decreased. From this reason, the upper limit of the effective Al concentration in the galvanizing bath is preferably set to be 0.500 mass %.

In order to improve the coating property and weldability, the surfaces of the hot-dip galvanized steel sheet and the alloyed hot-dip galvanized steel sheet of the present invention may be subjected to upper layer plating and to a variety of treatments, for example, a chromate treatment, a phosphate treatment, a lubricity-improving treatment, a weldability-improving treatment or the like.

After the hot-dip galvanizing and the alloying-heating treatment, the rolling is preferably performed for the purpose of the final shape straightening and the loss of yield-point elongation. When the elongation percentage is less than 0.2%, the effect is not sufficient. On the other hand, when the elongation percentage exceeds 1%, the yield ratio dramatically increases and the elongation deteriorates. Therefore, the elongation percentage is preferably set to be 0.2 to 1%. In addition, before dipping into the galvanizing bath, the steel sheet may be subjected to the pickling after the annealing to remove scales generated while annealing.

Examples

The effects of the present invention will be now described in more detail using Example. Incidentally, conditions of the examples are condition examples employed for confirming the applicability and effects of the present invention, and the present invention is not limited to these condition examples. The present invention can employ various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.

A hot-dip galvanized steel sheet and an alloyed hot-dip galvanized steel sheet were manufactured under conditions indicated in Tables 2 and 3 by casting steel having compositions indicated in Table 1. First, steel slabs having each composition were heated to a slab heating temperature (° C.) and then the hot rolling was finished at a finish rolling temperature (° C.). The obtained hot-rolled steel sheet was coiled at a coiling temperature (° C.) and then was subjected to the cold rolling indicated by a cold rolling ratio (%). Further, since a sheet shape is poor when the steel sheet was subjected to the cold rolling at a cold rolling ratio of 30% (cold rolling ratio less than 40%) and occurrence of scratches was concerned when the sheet subsequently passed through an annealing facility, the sheet-passing was abandoned. In addition, when the cold rolling was performed at a cold rolling ratio of 80% (cold rolling ratio exceeding 70%), the rolling load became a maximum value and a predetermined sheet thickness was not obtained, so that the cold rolling at the cold rolling ratio of 80% was abandoned.

After the cold rolling, the base steel sheet was annealed at the annealing temperature (° C.), was cooled (primary cooling) to a primary cooling temperature (° C.) at a primary cooling rate (° C./sec.), and then was cooled (secondary cooling) to a cooling stop temperature (° C.) at a secondary cooling rate (° C./sec.). Then, the base steel sheet was held at a holding temperature (° C.) for a holding time (sec.).

Thereafter, the base steel sheet was cooled to 70° C. or lower for the average cooling rate of 5° C./second or faster and then was rolled at a rolling ratio (elongation percentage) (%). The sheet thickness was 1.4 mm. Thereafter, the surface of the base steel sheet was polished and removed up to a depth of a surface polishing quantity (μm) and then was subjected to the Ni pre-plating with the Ni pre-plating weight (g/m²).

Next, the base steel sheet was heated to the heating temperature (° C.) at a temperature rising rate (° C./sec.) and then was subjected to the hot-dip galvanizing by dipping into the hot-dip galvanizing bath. In addition the base steel sheet was subjected to the alloying-heating treatment for an alloying time (sec.) at the alloying temperature (° C.), as necessary. Further, the speed of the jet flow (m/min) in the galvanizing bath was provided in a hot-dip galvanizing tank. In addition, the temperature of the galvanizing bath was managed to 440 to 470° C.

Thereafter, the base steel sheet was cooled to 70° C. or lower at the average cooling rate of 5° C./second or faster.

Mechanical properties, hole expandability (λ), a galvanized external appearance, an alloying degree, and plating adhesion of the obtained hot-dip galvanized steel sheet and the alloyed hot-dip galvanized steel sheet were estimated (Table 4). The mechanical properties were estimated by a tensile test based on JIS Z 2241. Tensile strength (TS) and total elongation (EL) were calculated from a stress-strain curve of the tensile test. Then, TS×EL and TS×λ as an indication of workability were calculated. The hole expandability was estimated by performing a hole expanding test based on Japan Iron and Steel Federation Standard JFS T 1001 to measure a hole expanding ratio. It was determined that the formability was good when the TS×EL was 17000 MPa·% or more and the TS×λ 40000 MPa·% or more. The galvanized external appearance was estimated as marks ◯ and X by determining the presence or absence of the galvanizing faults through visual observation. Here, the mark ◯ indicates that no galvanizing fault is present, and the mark X indicates that the galvanizing faults are present. The alloying Fe % indicates a mass % of Fe contained in the galvanized layer. In the alloyed hot-dip galvanized steel sheet subjected to the alloying treatment, the content of 7 to 15% Fe indicates that the alloying has well advanced. In the hot-dip galvanized steel sheet which is not subjected to the alloying treatment, the content of Fe may be 7% or less.

Experiments No. a, ba, bt, c, d, e, fa, g, ha, ht, i, j, k, l, m, n, and o are the present invention examples in which all of the properties pass and the steel sheet of the aimed property is obtained. On the other hand, in other Experiments No., in which composition or manufacturing method is out of the range of the present invention, several properties fail to pass.

TABLE 1

Steel type

No.
C
Si
Mn
P
S
N
Al
Al + Si
Others
Note

1
0.13
1.5
2.7
0.01
0.002
0.004
0.015
1.515
—
Inventive steel

2
0.18
1.5
2.4
0.012
0.003
0.0033
0.018
1.518
—
Inventive steel

3
0.19
2.7
1.6
0.011
0.004
0.0043
1.213
3.913
Cr: 0.96
Inventive steel

4
0.18
0.1
2.5
0.013
0.004
0.0022
0.51
0.61
Ce: 0.01, La: 0.002, V: 0.4
Inventive steel

5
0.28
1.8
1.6
0.012
0.01
0.0022
0.576
2.376
—
Inventive steel

6
0.27
0.6
2.4
0.01
0.0013
0.0024
1.751
2.351
Mg: 0.0008
Inventive steel

7
0.23
1.7
2.5
0.02
0.0023
0.0029
0.019
1.719
Ca: 0.008
Inventive steel

8
0.35
1.6
2.4
0.01
0.0014
0.0034
0.785
2.385
—
Inventive steel

9
0.19
1.5
1.9
0.02
0.002
0.0041
0.58
2.08
Ti: 0.03
Inventive steel

10
0.19
1.4
1.7
0.03
0.001
0.002
0.46
1.86
B: 0.001
Inventive steel

11
0.18
1.3
1.8
0.02
0.002
0.0024
0
1.3
Mo: 0.1
Inventive steel

12
0.2
1.8
2.6
0.03
0.001
0.0033
0.08
1.88
Cr: 0.8
Inventive steel

13
0.19
1.7
2.5
0.013
0.0015
0.0012
0.023
1.723
Nb: 0.051
Inventive steel

14
0.21
1.5
2.4
0.006
0.0042
0.0043
0.033
1.533
Ti: 0.056, B: 0.0053
Inventive steel

15
0.19
1.8
1.8
0.011
0.0032
0.0027
0.016
1.816
Mo: 0.33
Inventive steel

16

0.7
2.4
2.3
0.013
0.0047
0.0039
0.032
2.432
—

Comparative steel

17
0.18

3.4

2.5
0.014
0.0037
0.0015
0.048
3.448
—

Comparative steel

18
0.25
2

3.5

0.014
0.0049
0.0012
0.069
2.069
Ca: 0.015

Comparative steel

22
0.19
1.75
2.64
0.08
0.0015
0.005

2.36
4.11
Ca: 0.003

Comparative steel

23
0.35
0.1
2.4
0.08
0.0015
0.005
0.164

0.264

—

Comparative steel

24
0.09
2.87
1.8
0.0022
0.0006
0.002
0.02
2.89

Cr: 1.8

Comparative steel

25
0.11
1.7
2.9
0.0006
0.0007
0.0003
0.08
1.78

Ti: 0.9

Comparative steel

Underlines indicate that a numeral value is out of the range of the present invention

TABLE 2

Steel
Slab heating
Finish rolling
Coiling
Cold rolling
Annealing

Experiment
type
temperature
temperature
temperature
ratio
temperature

No.
No.
[° C.]
[° C.]
[° C.]
[%]
[° C.]

a
1
1230
900
500
50
880

ba
2
1230
890
640
55
820

bb
2
1220
890
650
55

700

bc
2
1200
900
590
55
790

bd
2
1180
920
620
55
800

be
2
1200
910
570
55
810

bf
2
1230
930
630
55
810

bg
2
1200
890
580
55
820

bh
2
1190
920
600
55
810

bi
2
1210
890
590
55
800

bj
2
1220
880
580
55
820

bk
2
1210
890
600
55
810

bl
2
1200
970
600
55
820

bm
2
1230
930
600
55
820

bn
2
1210
920
610
55
830

bo
2
1200
900
640
55
820

bp
2
1200
880
530
55
820

bq
2
1250
920
540
55
830

br
2
1240
920
580
55
820

bs
2
1230
900
600
55
820

bt
2
1230
920
580
55
820

c
3
1210
880
550
57
880

d
4
1230
930
500
49
850

e
5
1230
950
500
60
870

fa
6
1230
890
500
48
880

fb
6
1220
870
600
50

680

fc
6
1200
890
540
50
820

fd
6
1190
900
490
50
800

fe
6
1200
930
590
50
840

ff
6
1230
890
560
50
850

fg
6
1200
900
630
50
840

fh
6
1220
960
590
50
850

fi
6
1220
950
560
50
840

fj
6
1200
920
590
50
860

fk
6
1210
900
630
50
830

fl
6
1240
960
420
50
840

fm
6
1180
890
550
50
850

fn
6
1200
900
490
50
840

fo
6
1210
910
500
50
830

fp
6
1230
930
550
50
810

fq
6
1210
880
600
50
810

fr
6
1240
900
590
50
810

fs
6
1220
910
610
50
810

g
7
1250
900
600
49
830

ha
8
1250
980
550
50
880

hb
8
1220
910
560
50

650

hc
8
1180
890
600
50
840

hd
8
1200
900
620
50
830

he
8
1190
940
550
50
800

hf
8
1210
890
590
50
780

hg
8
1200
920
600
50
810

hh
8
1190
910
570
50
830

hi
8
1200
900
580
50
860

hj
8
1190
880
480
50
780

hk
8
1200
950
550
50
810

hl
8
1210
920
540
50
820

hm
8
1190
890
600
50
830

hn
8
1200
960
620
50
820

ho
8
1200
930
580
50
830

hp
8
1230
900
600
50
780

hq
8
1250
880
610
50
780

hr
8
1200
930
600
50
780

hs
8
1220
920
590
50
780

ht
8
1200
880
570
50
780

i
9
1230
920
630
50
840

j
10
1210
890
620
50
840

k
11
1230
870
650
45
850

l
12
1250
930
550
60
860

m
13
1220
940
530
50
840

n
14
1230
910
560
50
850

o
15
1240
950
600
50
860

p

16

1250
950
500
60
810

q

17

1230
960
600
50
780

r

18

1260
920
650
50
840

v

22

1200
890
450
50
850

w

23

1260
910
490
50
850

x

24

1240
900
420
50
850

y

25

1230
900
560
50
850

Primary cooling
Primary cooling
Secondary
Cooling stop

Experiment
temperature
rate
cooling rate
temperature

No.
[° C.]
[° C./sec]
[° C./sec]
[° C./sec]
Note

a
690
1.2
120
370
Present invention steel

ba
680
1.9
50
290
Present invention steel

bb
670
10.7
60
380

Comparative steel

bc

520

1.5
120
380

Comparative steel

bd
650
0.03
180
370

Comparative steel

be
670
1.2
8
400

Comparative steel

bf
740
1.8
60
370

Comparative steel

bg
710
1.2
50
350

Comparative steel

bh
730
1.6
60
420

Comparative steel

bi
740
1.5
50
400

Comparative steel

bj
710
2.7
70
380

Comparative steel

bk
720
2.3
80
390

Comparative steel

bl
710
3.8
120
390

Comparative steel

bm
710
3.8
100
400

Comparative steel

bn
680
3.8
120
390

Comparative steel

bo
680
4.5
60
420

Comparative steel

bp
680
1.2
80
70

Comparative steel

bq
690
1.2
80
50

Comparative steel

br
680
1.2
80

500

Comparative steel

bs
680
1.2
70

580

Comparative steel

bt
690
1.2
120
320
Present invention steel

c
690
3.4
80
360
Present invention steel

d
690
1.8
70
430
Present invention steel

e
690
2.9
60
400
Present invention steel

fa
720
13.5
50
390
Present invention steel

fb
680
2.4
100
400

Comparative steel

fc

530

2.4
120
400

Comparative steel

fd
650
0.06
140
410

Comparative steel

fe
670
1.6
6
390

Comparative steel

ff
720
2.4
80
430

Comparative steel

fg
700
1.6
60
450

Comparative steel

fh
720
2.2
50
370

Comparative steel

fi
710
1.6
70
390

Comparative steel

fj
690
2.8
100
400

Comparative steel

fk
720
2.4
130
400

Comparative steel

fl
700
5.6
120
390

Comparative steel

fm
690
2.4
90
390

Comparative steel

fn
710
6.2
80
390

Comparative steel

fo
710
2.4
60
400

Comparative steel

fp
700
2.4
70
70

Comparative steel

fq
680
2.4
70
50

Comparative steel

fr
680
2.4
60

520

Comparative steel

fs
680
2.4
60

590

Comparative steel

g
680
10.8
70
410
Present invention steel

ha
740
16.4
100
420
Present invention steel

hb
680
1.6
120
380

Comparative steel

hc

530

1.6
100
380

Comparative steel

hd
660
0.06
160
400

Comparative steel

he
690
1.4
4
400

Comparative steel

hf
730
1.6
80
380

Comparative steel

hg
680
1.8
50
370

Comparative steel

hh
710
2.4
60
410

Comparative steel

hi
720
2.8
190
400

Comparative steel

hj
740
2.4
100
390

Comparative steel

hk
700
2.6
80
390

Comparative steel

hl
710
4.8
60
380

Comparative steel

hm
700
5.8
70
360

Comparative steel

hn
730
2.6
70
430

Comparative steel

ho
720
1.6
80
400

Comparative steel

hp
680
1.6
80
70

Comparative steel

hq
680
1.6
80
50

Comparative steel

hr
690
1.2
60

520

Comparative steel

hs
680
1.2
60

600

Comparative steel

ht
660
2.1
80
200
Present invention steel

i
690
2.5
120
410
Present invention steel

j
680
2.4
90
400
Present invention steel

k
690
1.9
60
400
Present invention steel

l
690
10.9
80
390
Present invention steel

m
680
5.4
90
390
Present invention steel

n
680
8.2
100
380
Present invention steel

o
690
6.2
90
430
Present invention steel

p
730
1.9
60
430

Comparative steel

q
700
1.8
120
400

Comparative steel

r
650
2.5
70
420

Comparative steel

v
680
10.9
80
410

Comparative steel

w
630
11.3
80
390

Comparative steel

x
670
15.4
120
370

Comparative steel

y
620
13.5
100
380

Comparative steel

Underlines indicate that a numeral value is out of the range of the present invention

TABLE 3

Surface

Steel
Holding
Holding time at
Reduction
polishing
Ni pre-plating

Experiment
type
temperature
350 to 450° C.
ratio
quantity
weight

No.
No.
[° C.]
[sec]
[%]
[μm]
[g/m²]

a
1
370
150
0.4
0.15
0.21

ba
2
400
180
0.4
0.16
0.32

bb
2
380
178
0.6
0.35
1.4

bc
2
380
162
0.2
0.7
0.9

bd
2
370
250
0.3
0.15
0.4

be
2
400
420
0.4
0.21
0.3

bf
2
370
10
0.2
0.36
0.2

bg
2
350
240
0
0.21
0.2

bh
2
420
240

2.8

0.22
0.4

bi
2
400
120
0.3

0
1.3

bj
2
380
150
0.2
0.15
0.08

bk
2
390
240
0.3
0.12
0.8

bl
2
390
150
0.2
0.15
1.5

bm
2
400
150
0.3
0.15
1.5

bn
2
390
150
0.4
0.46
1.5

bo
2
420
150
0.3
0.46
1.5

bp
2
360
150
0.4
0.22
0.4

bq
2
90
150
0.4
0.24
0.4

br
2
450
150
0.4
0.18
0.4

bs
2

520

150
0.4
0.21
0.4

bt
2
380
220
0.3
0.24
2.24

c
3
360
170
0.4
0.17
0.26

d
4
430
120
0.2
0.26
1.24

e
5
400
190
0.3
0.13
1.87

fa
6
390
400
0.3
0.24
1.44

fb
6
400
500
0.4
0.22
0.4

fc
6
400
260
0.3
0.23
0.3

fd
6
410
500
0.4
0.16
0.6

fe
6
390
500
0.3
0.26
0.3

ff
6
430
0
0.3
0.2
0.4

fg
6
450
180
0
0.16
0.3

fh
6
370
200

1.6

0.12
0.4

fi
6
390
500
0.8

0.04

0.3

fj
6
400
250
0.8
0.16

0

fk
6
400
125
0.7
0.32
0.3

fl
6
390
240
0.6
0.18
0.2

fm
6
390
330
0.9
0.2
0.4

fn
6
390
500
0.7
0.16
0.5

fo
6
400
330
0.8
0.22
0.5

fp
6
370
200
0.2
0.43
0.22

fq
6
80
200
0.6
0.62
0.45

fr
6
450
200
0.4
0.26
0.25

fs
6

540

200
0.3
0.14
0.31

g
7
410
300
0.2
0.34
0.58

ha
8
420
260
0.2
0.16
0.24

hb
8
380
250
0.4
0.12
0.6

hc
8
380
260
0.3
0.18
0.3

hd
8
400
180
1
0.22
0.4

he
8
400
470
0.4
0.18
0.3

hf
8
380
0
0.9
0.24
0.3

hg
8
370
190
0
0.24
0.4

hh
8
410
320

2.2

0.16
0.3

hi
8
400
450
0.8

0.06

0.4

hj
8
390
260
0.7
0.18
0.05

hk
8
390
180
0.8
0.24
0.4

hl
8
380
220
0.5
0.16
0.8

hm
8
360
180
0.6
0.12
1.6

hn
8
430
460
0.8
0.16
0.8

ho
8
400
280
0.3
0.16
0.6

hp
8
350
330
0.4
0.22
0.4

hq
8
80
330
0.5
0.14
0.3

hr
8
450
330
0.4
0.32
0.2

hs
8

550

330
0.3
0.28
0.6

ht
8
380
640
0.4
0.62
0.2

i
9
410
180
0.4
0.15
0.32

j
10
400
190
0.2
0.19
0.22

k
11
400
250
0.3
0.26
0.99

l
12
390
280
0.5
0.24
1.5

m
13
390
800
0.4
0.18
0.23

n
14
380
260
0.3
0.11
0.33

o
15
430
300
0.6
0.06
0.27

p

16

430
120
0.3
0.16
1.55

q

17

400
130
0.4
0.24
1.24

r

18

420
250
0.5
0.52
1.5

v

22

410
250
0.2
0.12
0.24

w

23

390
260
0.7
0.12
0.26

x

24

370
200
0.6
0.15
0.35

y

25

380
120
0.4
0.15
0.45

Speed of jet

flow in

galvanizing
Temperature
Heating

Alloying

Experiment
bath
rising rate
temperature
Alloying time
temperature

No.
[m/min]
[° C./sec]
[° C.]
[sec]
[° C.]
Note

a
20
12.4
435
15
475
Present invention steel

ba
15
11.4
462
24
480
Present invention steel

bb
10
16.8
495
20
490

Comparative steel

bc
45
25.3
437
35
505

Comparative steel

bd
20
22.1
421
30
523

Comparative steel

be
20
24.5
460
25
510

Comparative steel

bf
25
26.5
468
31
545

Comparative steel

bg
25
25.3
478
24
490

Comparative steel

bh
20
28.4
469
19
500

Comparative steel

bi
20
22.2
458
25
510

Comparative steel

bj
20
24.6
423
15
520

Comparative steel

bk
4
17.6
485
20
500

Comparative steel

bl
20

3.2

427
19
540

Comparative steel

bm
25
20.5

390

15
520

Comparative steel

bn
20
22.4
456

130
520

Comparative steel

bo
20
10.5
450
15

620

Comparative steel

bp
15
12.4
450
20
510

Comparative steel

bq
15
20.3
460
20
500

Comparative steel

br
20
12.4
460
20
500

Comparative steel

bs
20
12.4
450
20
510

Comparative steel

bt
15
14.8
510
25
520
Present invention steel

c
20
20.6
486
35
520
Present invention steel

d
20
16.8
425
21
560
Present invention steel

e
15
10.4
437
13
590
Present invention steel

fa
25
16.5
480
14
490
Present invention steel

fb
30
10.8
480
20
480

Comparative steel

fc
15
16.8
470
16
490

Comparative steel

fd
10
16.8
460
24
480

Comparative steel

fe
15
16.8
450
30
500

Comparative steel

ff
15
10.5
480
16
500

Comparative steel

fg
20
16.8
480
24
490

Comparative steel

fh
15
20.4
450
30
500

Comparative steel

fi
25
12.4
460
8
540

Comparative steel

fj
20
10.9
450
16
510

Comparative steel

fk
7
20.4
460
30
490

Comparative steel

fl
20
1.8
480
24
490

Comparative steel

fm
45
16.8

370

30
480

Comparative steel

fn
20
20.5
490

200
540

Comparative steel

fo
20
30.7
500
28

650

Comparative steel

fp
20
10.6
520
28
500

Comparative steel

fq
20
20.8
510
28
500

Comparative steel

fr
20
25.4
500
28
490

Comparative steel

fs
20
20.4
500
24
500

Comparative steel

g
30
18.4
470
30
—
Present invention steel

ha
20
19.3
459
—
—
Present invention steel

hb
20
20.3
455
20
490

Comparative steel

hc
15
19.6
460
30
500

Comparative steel

hd
30
14.5
470
20
510

Comparative steel

he
25
20.5
480
28
520

Comparative steel

hf
25
10.8
455
24
490

Comparative steel

hg
20
15.6
480
20
480

Comparative steel

hh
15
16.8
500
24
490

Comparative steel

hi
30
30.8
460
45
500

Comparative steel

hj
40
20.4
460
20
510

Comparative steel

hk
5
18.6
480
30
480

Comparative steel

hl
20
0.8
500
20
470

Comparative steel

hm
50
18.2

290

24
480

Comparative steel

hn
25
20.8
460

200
510

Comparative steel

ho
10
26.8
480
30

630

Comparative steel

hp
15
10.8
470
30
480

Comparative steel

hq
20
20.5
450
24
480

Comparative steel

hr
15
14.6
470
45
480

Comparative steel

hs
20
14.9
460
30
480

Comparative steel

ht
30
16.8
460
30
510
Present invention steel

i
20
11.4
460
24
480
Present invention steel

j
20
25.5
450
35
490
Present invention steel

k
20
36.5
455
21
550
Present invention steel

l
20
19.7
449
—
—
Present invention steel

m
20
20.8
460
—
—
Present invention steel

n
20
8.8
450
16
530
Present invention steel

o
20
30.9
455
9
—
Present invention steel

p
40
10.6
462
15
480

Comparative steel

q
20
10.8
435
16
479

Comparative steel

r
20
12.5
458
17
510

Comparative steel

v
10
10.8
428
30
490

Comparative steel

w
15
15.5
437
25
475

Comparative steel

x
25
18.6
461
19
480

Comparative steel

y
30
14.6
478
—
—

Comparative steel

Underlines indicate that a numeral value is out of the range of the present invention

TABLE 4

Martensite
Martensite
Martensite
Bainite +
Residual

Experiment
Bainite
{circle around (1)}
{circle around (2)}
{circle around (3)}
Martensite
austenite
Ferrite
Pearlite

No.
[%]
[%]
[%]
[%]
[%]
[%]
[%]
[%]

a
47
14
3
0
64
8
28
0

ba
32
4
4
3
43
9
38
10

bb
0
0

0

0

0

0

100
0

bc
23
0

0

0

23

0

62

15

bd
23
0

0

0

23

0

55

22

be
38
0
0
0

38

0

49

13

bf
11
33
14
2
60

2

38
0

bg
35
10

0

0

45
16
39
0

bh
34
6
12
0
52

3

37
8

bi
40
6
4
3
53
8
39
0

bj
38
5
4
3
50
9
41
0

bk
40
5
4
3
52
10
38
0

bl
41
5
4
2
52
8
39
1

bm
43
4
5
3
55
9
36
0

bn
40
4
4
0
48

4

37

11

bo
38
5

0

0

43

0

39

18

bp
8
52

0

0

60

0

40
0

bq
7
55

0

0

62

2

36
0

br
33
5

0

0

38

0

39

23

bs
22
0

0

0

22

0

46

32

bt
29
10
7
6
52
9
39
0

c
53
6
3
3
65
11
24
0

d
55
8
4
0
67
10
23
0

e
48
8
7
3
66
13
21
0

fa
45
11
11
9
76
14
10
0

fb
0
0

0

0

0

0

100
0

fc
28
0

0

0

28

0

54

18

fd
24
0

0

0

24

0

49

27

fe
42
0
0
0
42

0

33

25

ff
10
65
3

0

78

4

18
0

fg
47
11

0

0

58
23
19
0

fh
45
11
27
0
83

2

15
0

fi
42
12
9
6
69
12
19
0

fj
43
11
9
6
69
13
18
0

fk
44
11
10
0
65
14
11
10

fl
41
12
11
6
70
11
19
0

fm
42
12
12
5
71
12
17
0

fn
41
13
8
6
68

0

16

16

fo
41
11

0

0

52

0

16

32

fp
11
58

0

0

69

3

28
0

fq
9
60

0

0

69

1

30
0

fr
39
0
0
0
39

0

38

23

fs
28
0
0
0
58

0

40

32

g
49
11
6
0
66
12
22
0

ha
52
13
10
7
82
18
0
0

hb
0
0

0

0

0
0
100
0

hc
26
0

0

0

26

0
37

37

hd
19
0

0

0

19

0
34

47

he
37
0
0
0
37
0
32

31

hf
0
55
12

0

67
15
18
0

hg
34
15

0

0

49
29
22
0

hh
52
14
29
0
95
5
0
0

hi
54
11
9
8
82
18
0
0

hj
42
11
8
6
67
15
18
0

hk
53
13
11
8
85
15
0
0

hl
52
12
10
8
82
18
0
0

hm
55
14
10
6
85
15
0
0

hn
62
13
11
9
95

5

0
0

ho
52
24

0

0

76

0

0

24

hp
12
61

0

0

73

3

24
0

hq
13
60

0

0

73

0

27
0

hr
35
7

0

0

42

0

39

19

hs
33
0

0

0

33

0

43

24

ht
20
21
6
4
51
21
28
0

i
52
4
4
3
63
10
27
0

j
51
9
4
0
64
10
26
0

k
54
4
3
4
65
9
26
0

l
50
8
4
4
66
8
26
0

m
55
5
3
4
67
9
24
0

n
44
8
6
5
63
10
27
0

o
51
6
7
7
71
9
20
0

p
0

63

22
0
85
15
0
0

q
16
5
5
3

29

10
61
0

r
10
75
13

0

98

2

0
0

v
16
0

0

0

16

13
71
0

w
92
8

0

0

100

0

0
0

x
62
21
11
3
97

3

0
0

y
5
0

0

0

5

0

95
0

Galvanized
Alloying

Experiment

external
Fe

No.
TS[MPa]
EL[%]
λ[% &]
TS × EL
TS × λ
appearance
[%]
Note

a
987
23
59
22701
58233
∘
7.8
Present invention steel

ba
1013
25
48
25325
48624
∘
10.4
Present invention steel

bb
683
21
67

14343

45761
∘
10.5

Comparative steel

bc
743
18
29

13374

21547

∘
11.1

Comparative steel

bd
789
15
25

11835

19725

∘
11.9

Comparative steel

be
824
16
29

13184

23896

∘
11.2

Comparative steel

bf
1139
11
24

12529

27336

∘
10.9

Comparative steel

bg
1037
26
24
26962

24888

∘
8.9

Comparative steel

bh
923
16
55

14768

50765
∘
9.3

Comparative steel

bi
984
25
45
24600
44280

x

11.3

Comparative steel

bj
989
23
43
22747
42527

x

11.5

Comparative steel

bk
982
23
51
22586
50082

x

9.4

Comparative steel

bl
981
21
41
20601
40221

x

12.3

Comparative steel

bm
989
23
44
22747
43516

x

11.9

Comparative steel

bn
923
14
51

12922

47073
∘
13.9

Comparative steel

bo
886
18
39

15948

34554

∘

16.3

Comparative steel

bp
1033
11
23

11363

23759

∘
10.7

Comparative steel

bq
1078
9
32
9702

34496

∘
10.2

Comparative steel

br
834
17
42

14178

35028

∘
9.9

Comparative steel

bs
753
16
46

12048

34638

∘
10.6

Comparative steel

bt
1001
22
65

22022

65065

∘
12.8
Present invention steel

c
1108
21
49
23268
54292
∘
11.6
Present invention steel

d
1239
20
45
24780
55755
∘
12.4
Present invention steel

e
1249
16
30
19984
37470
∘
8.7
Present invention steel

fa
1280
16
32
20480
40960
∘
7.9
Present invention steel

fb
712
20
72

14240

51264
∘
8.6

Comparative steel

fc
775
19
42

14725

32550

∘
8.8

Comparative steel

fd
793
17
38

13481

30134

∘
8.3

Comparative steel

fe
843
15
32

12645

26976

∘
9.5

Comparative steel

ff
1298
9
24

11682

31152

∘
9.7

Comparative steel

fg
1345
15
14
20175

18830

∘
8.9

Comparative steel

fh
1019
8
40
8152
40760
∘
11.3

Comparative steel

fi
1216
15
35
18240
42560

x

12.1

Comparative steel

fj
1224
14
40
17136
48960

x

11.4

Comparative steel

fk
1265
13
33
16445
41745

x

10.9

Comparative steel

fl
1274
16
34
20384
43316

x

9.6

Comparative steel

fm
1293
15
38
19395
49134

x

9.7

Comparative steel

fn
805
18
41

14490

33005

∘
14.7

Comparative steel

fo
785
19
42

14915

32970

∘

17.2

Comparative steel

fp
1352
7
9
9464

12168

∘
9.2

Comparative steel

fq
1397
6
4
8382
5588
∘
9.6

Comparative steel

fr
821
17
30

13957

24630

∘
9.6

Comparative steel

fs
782
18
24

14076

18768

∘
9.9

Comparative steel

g
1021
24
51
24504
52071
∘
2.8
Present invention steel

ha
1450
14
31
20300
44950
∘
2.7
Present invention steel

hb
759
20
59

15180

44781
∘
9.6

Comparative steel

hc
824
18
37

14832

30488

∘
10.2

Comparative steel

hd
793
16
32

12688

25376

∘
10.5

Comparative steel

he
857
13
30

11141

25710

∘
10.9

Comparative steel

hf
1652
4
7
6608

11564

∘
9.9

Comparative steel

hg
1523
13
7
19799

10661

∘
9.8

Comparative steel

hh
1221
10
36

12210

43956
∘
10.2

Comparative steel

hi
1434
13
32
18642
45888

x

11.4

Comparative steel

hj
1426
13
33
18538
47058

x

10.7

Comparative steel

hk
1462
14
34
20468
49708

x

10

Comparative steel

hl
1489
14
31
20846
46159

x

8.6

Comparative steel

hm
1482
13
38
19266
56316

x

8.1

Comparative steel

hn

1162

11
29

12782

33698

∘
13.8

Comparative steel

ho

1125

7
32
7875

36000

∘

16.8

Comparative steel

hp
1589
6
4
9534
6356
∘
10.1

Comparative steel

hq
1601
5
1
8005
1601
∘
9.7

Comparative steel

hr
923
16
28

14768

25844

∘
11.9

Comparative steel

hs
885
16
25

14160

22125

∘
9.9

Comparative steel

ht
1356
15
35
20340
47460
∘
10.3
Present invention steel

i
1185
17
45
20145
53325
∘
10.4
Present invention steel

j
1205
16
46
19280
55430
∘
11.6
Present invention steel

k
989
23
50
22747
49450
∘
12.4
Present invention steel

l
1201
17
35
20417
42035
∘
2.7
Present invention steel

m
1186
19
39
22534
46254
∘
1.6
Present invention steel

n
1208
17
34
20536
41072
∘
11.6
Present invention steel

o
1226
16
36
19616
44136
∘
1.7
Present invention steel

p
1950
5
3
9750
5850
∘
10.5

Comparative steel

q
945
18
26

17010

24570

x

6.8

Comparative steel

r
1523
9
29

13707

44167
∘
8.9

Comparative steel

v
894
17
24

15198

21456

∘
10.7

Comparative steel

w
927
11
62

10197

57474
∘
12.5

Comparative steel

x
862
18
24

15516

20688

∘
11.8

Comparative steel

y
1342
7
16
9394

21472

∘
2.1

Comparative steel

Underlines indicate that a numeral value is out of the range of the present invention

INDUSTRIAL APPLICABILITY

The present invention provides the high-strength galvanized steel sheet having excellent formability with the ultimate tensile strength of 980 MPa or more, which is suitable for the structural member, the reinforcing member, and the suspension member of automobiles. Accordingly, the present invention can be expected to greatly contribute to the lighter-weight of automobiles and is extremely high in effect in industry.

Number	Name	Date	Kind
8076008	Kaneko et al.	Dec 2011	B2
9783878	Wakabayashi	Oct 2017	B2
20080000555	Nonaka et al.	Jan 2008	A1
20120222781	Azuma et al.	Sep 2012	A1

Number	Date	Country
2781815	Jun 2011	CA
61-217529	Sep 1986	JP
5-59429	Mar 1993	JP
2526320	Jun 1996	JP
2526322	Jun 1996	JP
2005-105367	Apr 2005	JP
2006-283071	Oct 2006	JP
2007-270176	Oct 2007	JP
2009-242816	Oct 2009	JP
2010-65272	Mar 2010	JP
2010-209428	Sep 2010	JP
2011-111673	Jun 2011	JP
2011-168816	Sep 2011	JP
WO 2010103936	Sep 2010	WO
WO 2011065591	Jun 2011	WO

	Number	Date	Country
Parent	14347289		US
Child	15696650		US

High-strength hot-dip galvanized steel sheet and high-strength alloyed hot-dip galvanized steel sheet having excellent plating adhesion, formability, and hole expandability with tensile strength of 980 MPa or more and manufacturing method therefor

Information

Patent Number

Date Filed

Date Issued

Inventors

Original Assignees

Examiners

Agents

CPC

Field of Search

US

CPC

International Classifications

Term Extension

Abstract

Description

Claims

Priority Claims (1)

Parent Case Info

US Referenced Citations (4)

Foreign Referenced Citations (15)

Non-Patent Literature Citations (6)

Related Publications (1)

Divisions (1)

Entry
Brazilian Office Action for Appl. No. BR112014007496-8 dated May 9, 2018.
Brazilian Search Report for Appl. No. BR112014007496-8 dated May 9, 2018.
Indian Office Action issued in corresponding Indian Patent Application No. 3212/DELNP/2014 dated Nov. 30, 2018.
Canadian Office Action for Appl. No. 2,850,101 dated Jul. 7, 2015.
International Search Report for Appl. No. PCT/JP2012/075230 dated Dec. 25, 2012 (w/ English translation).
Korean Office Action for Appl. No. 10-2014-7010908 dated Aug. 11, 2015 (w/ English translation).