WO2018092735A1 - High strength steel sheet, production method therefor, and high strength galvanized steel sheet - Google Patents

Abstract

Provided is a high strength steel sheet that has a composition range including 0.001-0.200% of Ti and/or 0.001-0.200% of Nb among other components and has a steel structure wherein ferrite is present at an area ratio of 20% or more, the ferrite having an average crystalline grain size of 10-20 µm, bainite is present at an area ratio of 5% or more, martensite is present at an area ratio of 5% or more, the total area ratio of bainite and martensite is 15% or more, and residual austenite is present at a volume ratio of 5% or more. The high strength steel sheet has a TS of 780 MPa or more which is achieved by a texture of ferrite formed as a micro structure having an inverse intensity ratio of γ-fibers to α-fibers of more than 3.0, and has excellent ductibility and stiffness as well as deep drawability.

Description

High-strength steel sheet, method for producing the same, and high-strength galvanized steel sheet

The present invention relates to a high-strength steel sheet excellent in formability, which is mainly suitable as a structural member of an automobile, and a method for producing the same, and in particular, has a tensile strength (TS) of 780 MPa or more and is excellent not only in ductility but also in rigidity. Furthermore, the present invention seeks to obtain a high-strength steel sheet that is also excellent in deep drawability.
The high-strength steel sheet of the present invention includes a high-strength galvanized steel sheet having a galvanized layer on its surface.

In recent years, for the purpose of ensuring the safety of passengers in the event of a collision and improving fuel efficiency by reducing the weight of the vehicle body, there has been an aggressive movement to apply high-strength steel sheets with a thin plate thickness while maintaining TS of 780 MPa or more to automobile structural members. ing. In addition, recently, application of an extremely high strength high strength steel sheet having TS of 980 MPa class and 1180 MPa class has been studied.
However, in general, increasing the strength of a steel sheet causes a decrease in formability, so it is difficult to achieve both high strength and excellent formability, and a steel sheet having both high strength and excellent formability has been desired.

Also, recently, as a result of the remarkable progress in the strengthening of steel sheets, there is a movement to actively apply thin steel sheets having a TS of 780 MPa or more and a thickness of less than 2.0 mm. However, since the reduction in the rigidity of the vehicle body due to the thinning becomes a problem, it is necessary to improve the rigidity of the structural parts of the automobile. Since the rigidity of the structural part is determined by the plate thickness and Young's modulus of the steel sheet if the cross-sectional shape is the same, it is effective to increase the Young's modulus of the steel sheet in order to achieve both weight reduction and rigidity of the structural part.

The Young's modulus is largely governed by the texture of the steel sheet. In the case of iron having a body-centered cubic lattice, the Young's modulus is high in the <111> direction, which is the atomic dense direction, and conversely low in the <100> direction where the atomic density is small It has been known. It is known that the Young's modulus of normal iron with no anisotropy in crystal orientation is about 206 GPa, but by giving anisotropy to the crystal orientation and increasing the atomic density in a specific direction, Young's modulus can be increased. However, when considering the rigidity of an automobile body, since a load is applied from various directions, a steel sheet having a high Young's modulus in each direction is required in addition to a specific direction.

In response to such a request, for example, in Patent Document 1, in mass%, C: 0.02 to 0.15%, Si: 0.3% or less, Mn: 1.0 to 3.5%, P : 0.05% or less, S: 0.01% or less, Al: 1.0% or less, N: 0.01% or less, and Ti: 0.1 to 1.0%, the balance being Fe and inevitable A slab composed of mechanical impurities is hot-rolled, cold-rolled at a rolling reduction of 20 to 85%, and then recrystallized and annealed to have a ferrite single-phase microstructure, TS of 590 MPa or more, and the rolling direction. Proposed a method for producing a high-strength thin steel sheet with excellent rigidity, whose Young's modulus in the 90 ° direction is 230 GPa or more with respect to the rolling direction, and whose average Young's modulus in the 90 ° direction is 215 GPa or more with respect to the rolling direction. Has been.

Patent Document 2 discloses that in mass%, C: 0.05 to 0.15%, Si: 1.5% or less, Mn: 1.5 to 3.0%, P: 0.05% or less, S: 0.01% or less, Al: 0.5% or less, N: 0.01% or less, Nb: 0.02 to 0.15% and Ti: 0.01 to 0.15%, the balance A slab composed of Fe and inevitable impurities is hot-rolled, cold-rolled at a rolling reduction of 40 to 75%, and then recrystallized to have a mixed structure of ferrite and martensite, and TS is 590 MPa. As described above, a method for producing a high-rigidity and high-strength steel sheet having excellent workability and having a Young's modulus in a direction perpendicular to the rolling direction of 230 GPa or more has been proposed.

Further, in Patent Document 3, by mass%, C: 0.010 to 0.050%, Si: 1.0% or less, Mn: 1.0 to 3.0%, P: 0.005 to 0.00. 1%, S: 0.01% or less, Al: 0.005 to 0.5%, N: 0.01% or less and Nb: 0.03 to 0.3%, the balance being Fe and inevitable A slab made of impurities is cold-rolled after hot rolling, and recrystallized and annealed to have a steel structure including an area ratio of the ferrite phase of 50% or more and an area ratio of the martensite phase of 1% or more. A method for producing a high-strength steel sheet having a Young's modulus in the direction perpendicular to the rolling of 225 GPa or more and an average r value of 1.3 or more has been proposed.

JP 2007-092130 A JP 2008-240125 A JP 2005-120472 A

However, in the technique described in Patent Document 1, in order to achieve a tensile strength of 780 MPa or more, for example, referring to the examples, addition of V: 0.4 mass% and W: 0.5 mass% is necessary. is there. Moreover, in order to further increase the strength, it is indispensable to use expensive elements such as Cr and Mo, and there is a problem that the alloy cost increases.
In the technique described in Patent Document 2, it is effective to increase the Young's modulus in only one direction of the steel plate, but to improve the rigidity of the structural parts of automobiles that require a steel plate having a high Young's modulus in each direction. Not applicable.
The technology described in Patent Document 3 discloses that the rigidity and workability are excellent, and among the workability, it discloses that the deep drawability is particularly excellent, but TS is as low as about 660 MPa at most.
Further, the techniques described in Patent Documents 1 to 3 do not necessarily have the feature of being excellent not only in ductility and rigidity but also in deep drawability.

The present invention has been developed in view of such circumstances, and has developed a ferrite texture to γ-fiber (a fiber texture in which the <111> axis is parallel to the normal direction of the rolling surface) and utilizes bainite transformation. Then, by dispersing an appropriate amount of retained austenite, an object is to obtain a high-strength steel sheet having not only ductility but also rigidity and further excellent deep drawability while having a TS of 780 MPa or more.

In the present invention, excellent ductility, that is, El (total elongation) means that the value of TS × El is 15000 MPa ·% or more.
Also, excellent rigidity, that is, Young's modulus means that the Young's modulus in the rolling direction and the 45 ° direction with respect to the rolling direction is 205 GPa or more, and the Young's modulus in the direction perpendicular to the rolling direction is 220 GPa or more.
Furthermore, being excellent in deep drawability means that the average r value, which is an index of deep drawability, is 0.95 or more regardless of the strength of the steel sheet.

Now, the inventors have intensively studied to develop a high-strength steel sheet having a TS of 780 MPa or more and excellent in ductility, rigidity, and deep drawability.
As a result, the following knowledge was obtained.

(1) In the production method of the present invention, after adding steel of any one or two of Ti and Nb and heating the steel slab in which the composition of other alloy elements is appropriately controlled, In this case, the interstitial elements C and N are added due to the effect of promoting precipitation of the added Ti and / or Nb by making the coiling temperature (CT) of the hot rolling relatively high. It is important to deposit as carbides and nitrides with high thermal stability.
(2) After hot rolling, if necessary, heat treatment is performed to soften the hot-rolled sheet, and then the cold rolling is performed to increase the reduction ratio as much as possible to obtain α-fiber (<110> axis). It is important to develop a fiber texture that is parallel to the rolling direction) and a γ-fiber texture.
(3) The steel sheet structure before annealing obtained in this way is precipitated as solute C and solute N as carbides and nitrides with high thermal stability, and a texture of α-fiber and γ-fiber. When the ferrite is recrystallized by the first annealing treatment in the ferrite single phase region thereafter, the ferrite texture is changed to α-fiber and γ-fiber, particularly γ-fiber. As a result, the Young's modulus in all directions can be improved, and the average r value can be improved.
(4) Next, in the second annealing process in the ferrite + austenite two-phase region, a certain amount of austenite is generated while maintaining the texture of ferrite, and in the subsequent cooling process, bainite is generated to generate residual austenite. At the same time, it is possible to secure desired TS and El by generating ferrite, bainite and martensite in a certain ratio or more.
(5) Thus, it is possible to obtain a high-strength steel sheet that has a TS of 780 MPa or more and is excellent not only in ductility but also in rigidity and also in deep drawability.
The present invention has been made based on the above findings.

That is, the gist configuration of the present invention is as follows.
1. Ingredient composition is mass%,
C: 0.08% to 0.35%,
Si: 0.50% or more and 2.50% or less,
Mn: 1.50% or more and 3.00% or less,
P: 0.001% to 0.100%,
S: 0.0001% or more and 0.0200% or less and N: 0.0005% or more and 0.0100% or less,
Ti: 0.001% or more and 0.200% or less and Nb: 0.001% or more and 0.200% or less containing one or two selected from the balance, the balance consists of Fe and inevitable impurities,
Steel structure is area ratio,
Ferrite is 20% or more, the average grain size of the ferrite is 10 μm or more and 20 μm or less, and bainite is 5% or more,
Martensite is 5% or more,
The total area ratio of bainite and martensite is 15% or more,
In volume ratio, residual austenite is 5% or more,
Further, the ferrite texture has a microstructure in which the inverse strength ratio of γ-fiber to α-fiber is more than 3.0.
High strength steel plate.

2. In the high-strength steel plate according to 1 above, further, in mass%,
Al: 0.01% or more and 1.00% or less,
V: 0.005% or more and 0.100% or less,
B: 0.0001% to 0.0050%,
Cr: 0.05% or more and 1.00% or less,
Cu: 0.05% or more and 1.00% or less,
Sb: 0.0020% or more and 0.2000% or less,
Sn: 0.0020% or more and 0.2000% or less,
Ta: 0.0010% or more and 0.1000% or less,
Ca: 0.0003% or more and 0.0050% or less,
A high-strength steel sheet containing at least one element selected from Mg: 0.0003% to 0.0050% and REM: 0.0003% to 0.0050%.

3. The method for producing a high-strength steel sheet according to 1 or 2,
The steel slab having the component composition of 1 or 2 is heated to 1100 ° C. or higher and 1300 ° C. or lower, hot rolled at a finish rolling exit temperature: 800 ° C. or higher and 1000 ° C. or lower, and a coiling temperature: 300 ° C. or higher. Winding at 800 ° C. or lower, after pickling treatment, as it is or after holding in the temperature range of 450 ° C. or higher and 800 ° C. or lower for 900 s or more and 36000 s or less, cold rolling is performed at a rolling reduction of 40% or more,
Subsequently, the obtained cold-rolled sheet is heated to a temperature range of 450 ° C. or more and T1 temperature or less, and subjected to a first annealing treatment for holding at this temperature range for 300 s or more,
Next, after reheating to a temperature range of T1 temperature or more and T2 temperature or less and performing the second annealing treatment, cooling is performed at 300 ° C or more and 500 ° C or less at an average cooling rate of at least 550 ° C at 5 ° C / s or more. A method for producing a high-strength steel sheet, which is cooled to a stop temperature range and held for 10 s or longer in the cooling stop temperature range.
T1 temperature (° C.) = 720 + 29 × [% Si] -21 × [% Mn] + 17 × [% Cr]
T2 temperature (° C.) = 946−203 × [% C] ^1/2 + 45 × [% Si] −30 × [% Mn] + 150 × [% Al] −20 × [% Cu] + 11 × [% Cr] +350 × [% Ti] + 104 × [% V]
[% X] is the mass% of the component element X of the steel sheet, and zero for the component elements not contained.

4). A high-strength galvanized steel sheet having a galvanized layer on the surface of the high-strength steel sheet according to 1 or 2.

According to the present invention, it is possible to effectively obtain a high-strength steel sheet having a TS of 780 MPa or more and excellent not only in ductility but also in rigidity and also in deep drawability.
Therefore, by applying the high-strength steel plate obtained by the present invention to, for example, an automobile structural member, fuel efficiency can be improved by reducing the weight of the vehicle body, and the industrial utility value is extremely large.

Hereinafter, the present invention will be specifically described.
First, the reason why the component composition of the high-strength steel sheet is limited to the above range in the present invention will be described. In the following description, “%” representing the content of the constituent elements of steel means “mass%” unless otherwise specified.
[C: 0.08% to 0.35%]
C is an element indispensable for increasing the strength of a steel sheet and ensuring a stable amount of retained austenite, and is an element necessary for ensuring the amount of bainite and martensite and for retaining austenite at room temperature.
If the C content is less than 0.08%, it is difficult to secure desired TS and El. On the other hand, when the amount of C exceeds 0.35%, there is a concern of embrittlement or delayed fracture of the steel sheet, and the weldability and the heat-affected zone are remarkably hardened and the weldability deteriorates. Therefore, the C content is 0.08% or more and 0.35% or less. Preferably they are 0.12% or more and 0.33% or less, More preferably, they are 0.15% or more and 0.31% or less, More preferably, they are 0.20% or more and 0.30% or less.

[Si: 0.50% to 2.50%]
Si is an element useful for improving the El of the steel sheet by suppressing the formation of carbides and promoting the formation of retained austenite. It is also effective in suppressing the formation of carbides due to decomposition of retained austenite. Furthermore, since it has a high solid solution strengthening ability in ferrite, it contributes to improving the strength of steel. Further, Si dissolved in ferrite has an effect of improving work hardening ability and increasing the ductility of the ferrite itself.
In order to obtain such an effect, the Si amount needs to be 0.50% or more. On the other hand, if the amount of Si exceeds 2.50%, workability and toughness will deteriorate due to the increase in the amount of solid solution in ferrite, and surface properties will deteriorate due to the occurrence of red scale, etc. Causes degradation of plating adhesion and adhesion. Therefore, the Si amount is set to 0.50% or more and 2.50% or less. Preferably they are 0.80% or more and 2.00% or less, More preferably, they are 1.00% or more and 1.80% or less, More preferably, they are 1.10% or more and 1.70% or less.

[Mn: 1.50% to 3.00%]
Mn is effective for securing the strength of the steel sheet. In addition, the hardenability is improved to facilitate complex organization. At the same time, Mn acts to suppress the formation of pearlite during the cooling process, and facilitates transformation from austenite to bainite and martensite. In order to obtain such an effect, the amount of Mn needs to be 1.50% or more. On the other hand, if the amount of Mn exceeds 3.00%, Mn segregation in the thickness direction becomes remarkable, leading to a decrease in material stability. In addition, castability is deteriorated. Accordingly, the Mn content is 1.50% or more and 3.00% or less. Preferably they are 1.50% or more and 2.70% or less, More preferably, they are 1.80% or more and 2.50% or less.

[P: 0.001% to 0.100%]
P is an element that has a solid solution strengthening action and can be added according to a desired strength. In addition, it is an element effective for complex organization in order to promote ferrite transformation. In order to acquire such an effect, it is necessary to make P amount 0.001% or more. On the other hand, if the amount of P exceeds 0.100%, weldability is deteriorated and, when galvanizing is alloyed, the alloying speed is greatly delayed to impair the quality of galvanizing. Moreover, impact resistance is deteriorated by embrittlement due to grain boundary segregation. Therefore, the P amount is set to 0.001% or more and 0.100% or less. Preferably it is 0.005% or more and 0.050% or less.

[S: 0.0001% to 0.0200%]
S segregates at the grain boundaries and embrittles the steel during hot working, and also exists as a sulfide and reduces local deformability. Therefore, the steel content needs to be 0.0200% or less. On the other hand, the amount of S needs to be 0.0001% or more due to restrictions in production technology. Therefore, the S amount is set to 0.0001% or more and 0.0200% or less. Preferably it is 0.0001% or more and 0.0050% or less.

[N: 0.0005% to 0.0100%]
N is an element that greatly deteriorates the aging resistance of steel. In particular, when the amount of N exceeds 0.0100%, deterioration of aging resistance becomes remarkable, so the amount is preferably as small as possible. However, the amount of N needs to be 0.0005% or more due to restrictions on production technology. There is. Therefore, the N amount is set to 0.0005% or more and 0.0100% or less. Preferably it is 0.0005% or more and 0.0070% or less.

In the present invention, in addition to the above components, in order to obtain a ferrite having an orientation that is advantageous for improving the Young's modulus, Ti: 0.001% to 0.200% and Nb: 0.001% to 0.200 It is necessary to contain any 1 type or 2 types in% or less.
[Ti: 0.001% or more and 0.200% or less]
Ti forms precipitates with C, S, and N, and not only generates ferrite with an orientation that is advantageous for improving rigidity and deep drawability during annealing, but also suppresses coarsening of recrystallized grains. It also contributes to the improvement of strength. Moreover, when B is added, since N is precipitated as TiN, precipitation of BN is suppressed, and the effect of B described later is effectively expressed. In order to obtain such an effect, the Ti amount needs to be 0.001% or more. On the other hand, if the amount of Ti exceeds 0.200%, carbonitrides cannot be completely dissolved during normal slab reheating, and coarse carbonitrides remain, thereby increasing strength and suppressing recrystallization. The effect is not obtained. In addition, even when the continuously cast slab is cooled and then hot-rolled without being subjected to reheating, the contribution of the recrystallization suppression effect to the amount of Ti exceeding 0.200%. Is small, which increases the alloy cost. Therefore, the Ti amount is set to 0.001% or more and 0.200% or less. Preferably they are 0.030% or more and 0.170% or less, More preferably, they are 0.050% or more and 0.150% or less.

[Nb: 0.001% or more and 0.200% or less]
Nb forms fine precipitates at the time of hot rolling or annealing, and not only generates ferrite with an orientation that is advantageous for improving rigidity and deep drawability at the time of annealing, but also coarsens the recrystallized grains. Suppressing and contributing to the improvement of strength. In particular, when Nb is added in an appropriate amount, the austenite phase generated by reverse transformation at the time of annealing is refined, so that the microstructure after annealing is also refined to increase the strength. In order to obtain such an effect, the Nb amount needs to be 0.001% or more. On the other hand, if the Nb content exceeds 0.200%, the carbonitride cannot be completely dissolved at the time of normal slab reheating, and coarse carbonitrides remain, which increases the strength and suppresses recrystallization. The effect is not obtained. In addition, even when the continuously cast slab is cooled and then hot-rolled without being subjected to reheating, the contribution of the recrystallization suppressing effect to the amount of Nb exceeding 0.200%. Is small, which increases the alloy cost. Therefore, the Nb amount is set to be 0.001% or more and 0.200% or less. Preferably they are 0.030% or more and 0.170% or less, More preferably, they are 0.050% or more and 0.150% or less.

The high-strength steel sheet of the present invention is at least one selected from Al, V, B, Cr, Cu, Sb, Sn, Ta, Ca, Mg, and REM, if necessary, in addition to the above basic components. These elements can be contained alone or in combination. The balance of the component composition of the steel sheet is Fe and inevitable impurities.
[Al: 0.01% or more and 1.00% or less]
Al is an element effective for suppressing the formation of carbides and promoting the formation of retained austenite. Moreover, it is an element added as a deoxidizer in the steel making process. In order to obtain such effects, the Al amount needs to be 0.01% or more. On the other hand, when the Al content exceeds 1.00%, inclusions in the steel sheet increase and ductility deteriorates. Therefore, the Al content is set to 0.01% or more and 1.00% or less. Preferably they are 0.03% or more and 0.50% or less.

[V: 0.005% to 0.100%]
V forms fine precipitates during hot rolling or annealing, thereby suppressing coarsening of recrystallized grains, contributing to an increase in strength, and advantageous for improving rigidity and deep drawability during annealing. Produces a well-oriented ferrite. In order to obtain such an effect, the amount of V needs to be added by 0.005% or more. On the other hand, if the amount of V exceeds 0.100%, the moldability deteriorates. Therefore, when adding V, the content is made 0.005% or more and 0.100% or less.

[B: 0.0001% or more and 0.0050% or less]
B is an element effective for strengthening steel, and the effect of addition is obtained at 0.0001% or more. On the other hand, if B is added excessively over 0.0050%, the amount of martensite produced becomes excessive, and there is a concern that the ductility will decrease due to a significant increase in strength. Therefore, the B amount is set to 0.0001% or more and 0.0050% or less. Preferably it is 0.0005% or more and 0.0030% or less.

[Cr: 0.05% to 1.00%], [Cu: 0.05% to 1.00%]
Cr and Cu not only serve as solid solution strengthening elements, but also stabilize austenite and facilitate complex organization in the cooling process during annealing. In order to obtain such an effect, the Cr content and the Cu content must each be 0.05% or more. On the other hand, if the Cr content and the Cu content exceed 1.00%, the formability of the steel sheet is lowered. Therefore, when adding Cr and Cu, their contents are 0.05% or more and 1.00% or less, respectively.

[Sb: 0.0020% or more and 0.2000% or less], [Sn: 0.0020% or more and 0.2000% or less]
Sb and Sn are added as necessary from the viewpoint of suppressing decarburization in the region of several tens of μm of the steel sheet surface layer caused by nitriding and oxidation of the steel sheet surface. This is because suppressing such nitriding and oxidation prevents the martensite generation amount on the steel sheet surface from decreasing and is effective in ensuring the strength and material stability of the steel sheet. On the other hand, if any of these elements is added excessively exceeding 0.2000%, the toughness is reduced. Accordingly, when Sb and Sn are added, their contents are within the range of 0.0020% or more and 0.2000% or less, respectively.

[Ta: 0.0010% to 0.1000%]
Ta, like Ti and Nb, generates carbides and carbonitrides and contributes to high strength. In addition, a part of the Nb carbide or Nb carbonitride is solid-solved to form a composite precipitate such as (Nb, Ta) (C, N), which significantly suppresses the coarsening of the precipitate. It is thought that it contributes effectively to the strength improvement of a steel plate by the object stabilization effect. Therefore, it is preferable to contain Ta.
Here, the above-mentioned precipitate stabilization effect can be obtained by setting the Ta content to 0.0010% or more. On the other hand, even if Ta is added excessively, the precipitate stabilization effect is saturated. , Alloy costs increase. Therefore, when Ta is added, the content is within the range of 0.0010% to 0.1000%.

[Ca: 0.0003% or more and 0.0050% or less], [Mg: 0.0003% or more and 0.0050% or less], [REM: 0.0003% or more and 0.0050% or less]
Ca, Mg, and REM are elements used for deoxidation, and are effective elements for spheroidizing the shape of the sulfide and improving the adverse effect of the sulfide on local ductility. In order to obtain these effects, 0.0003% or more must be added. However, when Ca, Mg and REM are added in excess of 0.0050%, inclusions and the like are increased to cause defects on the surface and inside. Therefore, when Ca, Mg and REM are added, their contents are 0.0003% or more and 0.0050% or less, respectively.

Next, the microstructure (steel structure) of the high-strength steel sheet of the present invention will be described.
[Area ratio of ferrite: 20% or more, average grain size of ferrite: 10 μm or more and 20 μm or less]
In the present invention, this is a very important invention constituent element. The high-strength steel sheet of the present invention comprises a composite structure in which retained austenite responsible for ductility and bainite and martensite responsible for strength are dispersed in soft ferrite that contributes to higher Young's modulus and higher r-value. In order to ensure a sufficient ductility and strength-ductility balance, the area ratio of ferrite generated in the second annealing and cooling process needs to be 20% or more.
The upper limit of the area ratio of ferrite is not particularly limited, but is preferably 80% or less, more preferably 70% or less, and still more preferably 60% or less for securing the strength.
In addition, it is important that the average crystal grain size of ferrite is 10 μm or more and 20 μm or less. If the average crystal grain size of ferrite is less than 10 μm, a desired texture cannot be obtained, and a desired Young's modulus and average r value cannot be ensured. On the other hand, if the average crystal grain size of ferrite exceeds 20 μm, the desired area ratio of bainite and martensite cannot be obtained, and the desired strength cannot be ensured. Therefore, the average crystal grain size of ferrite is 10 μm or more and 20 μm or less. Preferably they are 11 micrometers or more and 17 micrometers or less. Such control of the ferrite grain size can be achieved by appropriately controlling the annealing temperature and the holding time in the manufacturing process, particularly in the first annealing treatment.

[Bainite area ratio: 5% or more]
In the present invention, this is a very important invention constituent element. The formation of bainite is necessary for concentrating C in untransformed austenite and obtaining retained austenite that can exhibit the TRIP effect in a high strain region during processing. In order to achieve both high strength and high ductility, it is effective to increase the amount of retained austenite produced.
In the present invention, in the holding process after the second annealing, when the amount of bainite produced is less than 5%, C concentration to austenite does not sufficiently proceed, so that a residual that exhibits a TRIP effect in a high strain region during processing. The amount of austenite decreases. Moreover, since the fraction of untransformed austenite in the holding process after the second annealing is increased and the fraction of martensite after cooling is increased, TS is increased, but ductility is decreased. Therefore, the area ratio of bainite needs to be 5% or more in terms of the area ratio with respect to the entire steel sheet structure.
The upper limit of the area ratio of bainite is not particularly limited, but is preferably 60% or less, more preferably 50%, in order to ensure the area ratio of ferrite that is advantageous for increasing the Young's modulus and the r value. % Or less.

[Martensite area ratio: 5% or more]
In the present invention, the martensite area ratio needs to be 5% or more in order to ensure the strength of the steel sheet.
The upper limit of the martensite area ratio is not particularly limited, but in order to ensure a good ductility at the same time as securing an area ratio of ferrite advantageous for increasing the Young's modulus and r value, martensite The area ratio is preferably 50% or less, more preferably 40% or less.

[Area ratio of bainite and martensite: 15% or more in total]
In the present invention, this is a very important invention constituent element. When the area ratio of bainite and martensite, which are low-temperature transformation phases, is less than 15%, the strength of the steel sheet cannot be ensured. Therefore, in order to secure a desired TS, the area ratio of bainite and martensite needs to be 15% or more in total.
The upper limit of the total area ratio of bainite and martensite is not particularly limited, but is preferably 70% or less for securing the area ratio of ferrite that is advantageous for increasing the Young's modulus and increasing the r value. More preferably, it is 60% or less, More preferably, it is 55% or less.

The area ratio of ferrite, bainite, and martensite is 1 vol. After polishing the plate thickness section (L section) parallel to the rolling direction of the steel sheet. Corrosion with% nital, and a plate thickness of 1/4 position (position corresponding to 1/4 of the plate thickness in the depth direction from the steel plate surface) using a scanning electron microscope (SEM) at a magnification of 3000 times The area ratio of the constituent phases (ferrite, bainite and martensite) was calculated for the three visual fields using Adobe Photoshop from Adobe Systems, and the values were averaged. Can be sought. In the above structure image, ferrite has a gray structure (base structure), martensite has a white structure, and bainite has a structure in which a white structure is mixed with a gray base.
Further, the average crystal grain size of ferrite can be obtained as follows. Using the above-mentioned Adobe Photoshop, the value obtained by correcting the length of the line segment drawn on the image to the actual length was divided by the number of crystal grains passing through the line segment drawn on the image.

[Volume ratio of retained austenite: 5% or more]
In the present invention, in order to ensure a good ductility and strength-ductility balance, the amount of retained austenite needs to be 5% or more by volume ratio. In order to secure a better ductility and strength-ductility balance, the amount of retained austenite is preferably 8% or more, more preferably 11% or more in terms of volume ratio. Here, the upper limit of the volume ratio of retained austenite is not particularly limited, but is preferably 20% or less.
The volume fraction of retained austenite was determined by measuring the X-ray diffraction intensity after grinding and polishing the steel plate to ¼ of the plate thickness in the plate thickness direction. For incident X-rays, Co—Kα is used, and the amount of retained austenite is calculated from the intensity ratio of each surface of (200), (220), (311) of austenite to the diffraction intensity of each surface of (200), (211) of ferrite. Calculated.

Further, in the microstructure according to the present invention, in addition to the above-described ferrite, bainite, martensite, and retained austenite, carbides such as tempered martensite, tempered bainite, pearlite, cementite, and other known structures may be included. However, the effect of the present invention is not impaired even if these total amounts are within the range of 15% or less in terms of area ratio.

Next, the texture of the steel plate will be described.
[Inverse strength ratio of γ-fiber to α-fiber of ferrite texture: more than 3.0]
α-fiber is a fiber texture whose <110> axis is parallel to the rolling direction, and γ-fiber is a fiber texture whose <111> axis is parallel to the normal direction of the rolling surface. The body-centered cubic metal is characterized in that α-fiber and γ-fiber are strongly developed by rolling deformation and a texture belonging to them is formed even by recrystallization annealing.
In the present invention, when the inverse strength ratio of γ-fiber to α-fiber in the ferrite texture is 3.0 or less, the degree of integration of γ-fiber suitable for high Young's modulus and high r-value is low. It is difficult to ensure the Young's modulus and the average r value. Therefore, the inverse strength ratio of γ-fiber to α-fiber in the ferrite texture needs to exceed 3.0. The upper limit of the inverse intensity ratio is not particularly limited, but is preferably 8.0 or less.
In the high-strength steel sheet obtained by the conventional general manufacturing method, the inverse strength ratio of γ-fiber to α-fiber is about 1.0 to 2.5.

The inverse strength ratio of γ-fiber to α-fiber in the ferrite texture was smoothed by wet polishing and buffing using a colloidal silica solution on the plate thickness section (L section) parallel to the rolling direction of the steel sheet. Thereafter, 0.1 vol. Corrosion with% nital reduces asperities on the sample surface as much as possible, and completely removes the work-affected layer, and then corresponds to 1/4 position of the plate thickness (1/4 of the plate thickness in the depth direction from the steel plate surface). The crystal orientation is measured using SEM-EBSD (Electron Back-Scatter Diffraction). Next, using the OIM Analysis of AMETEK EDAX, first select bainite and martensite containing adjacent ferrites of similar orientation using the highlight grain function, and then select the bainite and martensite using the chart function. By extracting only azimuth information, it can be calculated by separating the texture information of ferrite and bainite and martensite, and determining the inverse strength of α-fiber and γ-fiber of the separated ferrite.

Next, the manufacturing method of the high strength steel plate of this invention is demonstrated.
In the present invention, any one or two elements of Ti and Nb are added, and the steel slab in which the composition of other alloy elements is appropriately controlled is heated and subjected to hot rolling. At that time, the hot rolling coiling temperature (CT) is set to a relatively high temperature, so that the interstitial elements C and N are highly thermally stable due to the precipitation promoting effect of the added Ti and / or Nb. It is important to deposit as carbides and nitrides.
Moreover, after hot rolling, if necessary, heat treatment is performed to soften the hot-rolled sheet. Thereafter, when cold rolling is performed, it is important to develop the texture of α-fiber and γ-fiber by increasing the reduction ratio as much as possible.
As described above, the steel sheet structure before the annealing treatment is a structure in which solute C and N are precipitated as carbides and nitrides having high thermal stability, and a texture of α-fiber and γ-fiber is developed. Then, by recrystallizing the ferrite in the first heating step (first annealing treatment) in the subsequent ferrite single phase region, the texture of the ferrite is developed into α-fiber and γ-fiber, particularly γ-fiber. As a result, the Young's modulus in all directions can be improved, and the average r value can be improved.
Next, in the second heating step (second annealing process) in the ferrite + austenite two-phase region, a certain amount of austenite is generated while maintaining the ferrite texture, and bainite and residual austenite are generated in the subsequent cooling process. At the same time, by producing ferrite, bainite and martensite at a certain ratio or more, a high strength steel sheet having a TS of 780 MPa or more, excellent in not only ductility but also rigidity, and also excellent in deep drawability is obtained. Is possible.
Moreover, the high-strength galvanized steel sheet of the present invention can be manufactured by subjecting the above-described high-strength steel sheet to a publicly known galvanizing treatment.

Hereinafter, each manufacturing process will be described.
[Heating temperature of steel slab: 1100 ° C or higher and 1300 ° C or lower]
Precipitates present in the heating stage of the steel slab exist as coarse precipitates in the finally obtained steel sheet and do not contribute to strength, so the Ti and Nb-based precipitates precipitated during casting are redissolved. There is a need.
Here, when the heating temperature of the steel slab is less than 1100 ° C., not only is it difficult to sufficiently dissolve the precipitate, but there is a problem that the risk of trouble occurring during hot rolling due to an increase in rolling load increases. . In addition, it is necessary to scale off defects such as bubbles and segregation in the surface layer of the slab, reduce cracks and irregularities on the steel sheet surface, and achieve a smooth steel sheet surface. Furthermore, when the precipitate produced | generated at the time of casting does not melt | dissolve but remains as a coarse precipitate, the problem that El, Young's modulus, and average r value fall also arises. Furthermore, there is a concern that residual austenite cannot be produced effectively and ductility is lowered. Therefore, the heating temperature of the steel slab of the present invention needs to be 1100 ° C. or higher. On the other hand, when the heating temperature of the steel slab exceeds 1300 ° C., the scale loss increases as the oxidation amount increases. Therefore, the heating temperature of the steel slab needs to be 1300 ° C. or lower.
Therefore, the heating temperature of the steel slab is set to 1100 ° C. or higher and 1300 ° C. or lower. Preferably they are 1150 degreeC or more and 1280 degrees C or less, More preferably, they are 1150 degreeC or more and 1250 degrees C or less.

[Finishing rolling delivery temperature: 800 ° C or higher and 1000 ° C or lower]
The heated steel slab is hot-rolled by rough rolling and finish rolling to form a hot-rolled steel sheet. At this time, when the finish rolling exit temperature exceeds 1000 ° C., the amount of oxide (scale) generated increases rapidly, the interface between the base iron and the oxide becomes rough, and the surface quality after pickling and cold rolling is high. It tends to deteriorate. In addition, if there is a part of the hot rolled scale remaining after pickling, the ductility is adversely affected. Furthermore, the crystal grain size becomes excessively coarse, and the surface of the pressed product may be roughened during processing.
On the other hand, if the finish rolling exit temperature is less than 800 ° C., the rolling load increases and the rolling load increases. In addition, the reduction ratio of austenite in an unrecrystallized state increases, abnormal texture develops, in-plane anisotropy in the final product becomes remarkable, and material uniformity and material stability are not only impaired. , El, Young's modulus, and average r value itself also decrease.
Therefore, the finish rolling exit temperature of hot rolling needs to be 800 ° C. or higher and 1000 ° C. or lower. Preferably they are 820 degreeC or more and 950 degrees C or less.

Incidentally, the steel slab is preferably manufactured by a continuous casting method in order to prevent macro segregation, but it can also be manufactured by an ingot-making method or a thin slab casting method. In addition to the conventional method in which the steel slab is manufactured and then cooled to room temperature and then heated again, the steel slab is not cooled to room temperature. Energy-saving processes such as direct feed rolling and direct rolling that are rolled immediately after application can also be applied without problems. The slab is made into a sheet bar by rough rolling under normal conditions. However, if the heating temperature is lowered, the sheet is heated using a bar heater before finishing rolling in order to prevent problems during hot rolling. It is preferred to heat the bar.

[Winding temperature after hot rolling: 300 ° C or higher and 800 ° C or lower]
When the coiling temperature after hot rolling exceeds 800 ° C, the crystal grain size of ferrite in the hot-rolled sheet structure increases, and the carbonitrides of Ti and Nb become coarse, so that during cold rolling and annealing Of γ-fiber becomes weak, and it becomes difficult to secure a desired Young's modulus and average r value. On the other hand, when the coiling temperature after hot rolling is less than 300 ° C., the hot rolled sheet strength increases, the rolling load in cold rolling increases, and the productivity decreases. Also, if cold rolling is performed on a hard hot-rolled sheet mainly composed of martensite, minute internal cracks (brittle cracks) are likely to occur along the former austenite grain boundaries of martensite, and the ductility of the final annealed sheet decreases. To do. Therefore, the coiling temperature after hot rolling needs to be 300 ° C. or higher and 800 ° C. or lower. Preferably they are 350 degreeC or more and 700 degrees C or less, More preferably, they are 380 degreeC or more and 650 degrees C or less.
Note that rough rolling sheets may be joined to each other during hot rolling to continuously perform finish rolling. Moreover, you may wind up a rough rolling board once. Moreover, in order to reduce the rolling load during hot rolling, part or all of the finish rolling may be lubricated rolling. Performing lubrication rolling is also effective from the viewpoint of uniform steel plate shape and uniform material. In addition, it is preferable to make the friction coefficient at the time of lubrication rolling into the range of 0.10 or more and 0.25 or less.

The hot-rolled steel sheet thus manufactured is pickled. Since pickling can remove oxides on the surface of the steel sheet, it is important for ensuring good chemical conversion properties and plating quality in the final high-strength steel sheet. Moreover, pickling may be performed once or may be divided into a plurality of times.
After the above pickling treatment, it is kept as it is or in a temperature range of 450 ° C. to 800 ° C. for a time of 900 s to 36000 s, and then cold-rolled at a reduction ratio of 40% or more.

[Heat treatment temperature range and holding time after hot-rolled plate pickling treatment: Hold for 900 s to 36000 s in a temperature range of 450 ° C. to 800 ° C.]
When the heat treatment temperature range is less than 450 ° C. or the heat treatment holding time is less than 900 s, tempering after hot rolling is insufficient, and at least one of ferrite, pearlite, bainite and martensite is mixed during the subsequent cold rolling. The resulting structure becomes uneven, and the uniform refinement becomes insufficient under the influence of the hot-rolled sheet structure. As a result, the ratio of the coarse low temperature transformation phase increases in the structure of the final annealed plate, resulting in a non-uniform structure, and the El, Young's modulus, and average r value of the final annealed plate may decrease.
On the other hand, when the heat treatment holding time exceeds 36000 s, productivity may be adversely affected. In addition, when the heat treatment temperature range exceeds 800 ° C., it becomes a non-uniform and hardened coarse two-phase structure of ferrite and martensite or pearlite, and becomes a non-uniform structure before cold rolling. The ratio of coarse martensite may increase, and the El, Young's modulus, and average r value of the final annealed sheet may also decrease.
Therefore, when heat treatment is performed after the hot-rolled sheet pickling treatment, the temperature range needs to be 450 ° C. or higher and 800 ° C. or lower, and the holding time needs to be 900 s or higher and 36000 s or lower.

[Draft ratio during cold rolling: 40% or more]
Cold rolling is performed after the hot rolling step to accumulate α-fiber and γ-fiber effective in improving Young's modulus and average r value. That is, by developing α-fiber and γ-fiber by cold rolling, the ferrite having α-fiber and γ-fiber, especially γ-fiber, is increased in the structure after the subsequent annealing process, and Young's modulus and average Increase the r value. In order to obtain such an effect, the rolling reduction during cold rolling needs to be 40% or more. Further, from the viewpoint of improving the Young's modulus and the average r value, the rolling reduction is preferably 45% or more, more preferably 50% or more. In addition, about the frequency | count of a rolling pass and the rolling reduction for every pass, the effect of this invention can be acquired, without being specifically limited. Moreover, although there is no limitation in particular in the upper limit of the said rolling reduction, it is about 80% industrially.

[Temperature range of first annealing treatment: 450 ° C. or higher and T1 temperature or lower]
In the present invention, this is a very important invention constituent element. When the annealing temperature range of the first firing treatment is less than 450 ° C., a large amount of unrecrystallized ferrite remains, and the amount of ferrite having γ-fiber formed during recrystallization of ferrite decreases, and the Young's modulus and average r in each direction The value drops. On the other hand, when the temperature of the first annealing treatment exceeds the T1 temperature, austenite is first nucleated from the nucleation site of recrystallized ferrite having γ-fiber, which is suitable for improvement of Young's modulus and average r value. The area ratio of ferrite having γ-fiber is reduced. In addition, the volume fraction of austenite generated during annealing increases, and the volume fraction of ferrite accumulated in α-fiber and γ-fiber, especially γ-fiber, decreases, resulting in a decrease in Young's modulus and average r value in each direction. To do. Furthermore, when carrying out the cooling step after heating, ferrite, martensite, tempered martensite, bainite, tempered bainite, or carbides such as pearlite, cementite, etc. that are generated by transformation of austenite during cooling increase, γ Since the area ratio of ferrite accumulated in -fiber decreases, it becomes difficult to accumulate in α-fiber and γ-fiber, particularly γ-fiber. Therefore, the temperature range of the first annealing process needs to be 450 ° C. or higher and T1 temperature or lower. Furthermore, from the viewpoint of improving the Young's modulus and the average r value, the temperature range of the first annealing treatment is preferably 500 ° C. or higher and T1 temperature or lower, more preferably 550 ° C. or higher and T1 temperature or lower. Here, the T1 temperature means the Ac ₁ point.

[Retention time in the first annealing treatment: 300 s or more]
In the present invention, this is a very important invention constituent element. When the holding time in the first annealing treatment is less than 300 s, unrecrystallized ferrite remains, and the degree of accumulation in γ-fiber decreases, so that the Young's modulus and average r value in each direction decrease. For this reason, holding time shall be 300 s or more. Further, there is no particular limitation, but if the holding time exceeds 100,000 s, the recrystallized ferrite grain size becomes coarse and it becomes difficult to secure a desired TS. Therefore, the holding time is preferably 100,000 s or less. . Accordingly, the holding time is 300 s or longer. Preferably it is 300 s or more and 100,000 or less, More preferably, it is 300 or more and 36000 s or less, More preferably, it is 300 or more and 21600 s or less.

The heat treatment method may be any of continuous annealing and batch annealing. Moreover, when implementing a cooling process after the 1st annealing process, you may cool to room temperature and you may perform the process which passes an overaging zone. The cooling method and cooling rate in the cooling step are not particularly defined, and any cooling such as furnace cooling in batch annealing, air cooling, and gas jet cooling, mist cooling, and water cooling in continuous annealing may be used. The pickling may be performed according to a conventional method. Although there is no particular limitation, since the steel sheet shape may be deteriorated when the average cooling rate to room temperature or overaging zone exceeds 80 ° C./s, the average cooling rate is 80 ° C./s or less. It is preferable to do.

[Temperature range of second annealing treatment: T1 temperature or more and T2 temperature or less]
When the second annealing temperature is lower than the T1 temperature, austenite formation is insufficient during annealing, and as a result, a sufficient amount of low-temperature transformation phase is obtained in the cooling and subsequent holding processes after the second annealing treatment. Therefore, it becomes difficult to secure a desired TS. Further, since the bainite transformation during holding is delayed, a sufficient amount of retained austenite cannot be secured, and it becomes difficult to improve ductility. On the other hand, when the annealing temperature of the second time exceeds the T2 temperature, it becomes an austenite single-phase temperature range, so that the ferrite texture formed in the heating step and the holding step after heating is randomized and finally obtained. The Young's modulus and average r value of the steel sheet are lowered. Moreover, since the area ratio of the ferrite which contributes to ductility improvement also falls, it becomes difficult to ensure desired El. Therefore, the temperature range of the second annealing treatment is set to T1 temperature or more and T2 temperature or less. The holding time of the second annealing treatment is not particularly limited, but is preferably 10 s or more and 1000 s or less. Here, the T2 temperature means the Ac ₃ point.

[Average cooling rate to 550 ° C. after second annealing treatment: 5 ° C./s or more]
In the cooling step after the second annealing treatment, if the average cooling rate to at least 550 ° C. is less than 5 ° C./s, the untransformed austenite is transformed into pearlite, and a desired amount of bainite and martensite cannot be secured, and the desired It becomes difficult to secure TS and El. Moreover, although it is not necessary to specifically limit, since the above-described average cooling rate exceeds 200 ° C./s, the shape of the steel sheet may be deteriorated and it may be difficult to control the cooling reaching temperature. Is preferably 200 ° C./s or less. Therefore, the average cooling rate to at least 550 ° C. in the cooling step after reheating is set to 5 ° C./s or more. Preferably they are 5 degreeC / s or more and 200 degrees C / s or less, More preferably, they are 8 degreeC / s or more and 80 degrees C / s or less, More preferably, they are 10 degreeC / s or more and 50 degrees C / s or less.

[Cooling stop temperature after second annealing treatment: 300 ° C. or more and 500 ° C. or less]
If the cooling stop temperature after the second annealing treatment is less than 300 ° C., the untransformed austenite is transformed into only martensite when cooling is stopped after the second annealing treatment, so that the desired amount of bainite cannot be secured. The desired El cannot be secured. On the other hand, if the cooling stop temperature after the second annealing treatment exceeds 500 ° C., untransformed austenite is transformed into pearlite, and a desired amount of bainite and martensite cannot be secured, and it is difficult to secure desired TS and El. It becomes. Therefore, the cooling stop temperature after the second annealing treatment is set to 300 ° C. or more and 500 ° C. or less. Furthermore, from the viewpoint of improving the balance between strength and ductility, the cooling stop temperature after the second annealing treatment is preferably set to 300 ° C. or higher and 480 ° C. or lower. More preferably, it is 350 degreeC or more and 460 degreeC or less.

[Holding time in cooling stop temperature range: 10s or more]
When the holding time in the reheating temperature region is less than 10 s, the time for C concentration to austenite to proceed becomes insufficient, and it becomes difficult to finally secure a desired volume ratio of retained austenite. Although it is not necessary to specifically limit, when it stays for more than 1000 s, the volume ratio of retained austenite does not increase, and a remarkable improvement in ductility is not confirmed, so that it tends to be saturated. Accordingly, the holding time in the cooling stop temperature region is set to 10 s or more. Preferably, it is 10 seconds or more and 1000 seconds or less.
The cooling after the holding does not need to be specified, and may be cooled to a desired temperature by any method. The desired temperature is preferably about room temperature.

[Zinc plating treatment]
When hot dip galvanizing treatment is performed, the steel plate subjected to the annealing treatment is immersed in a galvanizing bath at 440 ° C. or higher and 500 ° C. or lower to perform hot dip galvanizing treatment, followed by gas wiping etc. Adjust. For hot dip galvanization, it is preferable to use a galvanizing bath having an Al content of 0.10 mass% or more and 0.23 mass% or less. In addition, when the galvanizing alloying treatment is performed, the galvanizing alloying treatment is performed in the temperature range of 470 ° C. or more and 600 ° C. or less after the hot dip galvanizing treatment. When the alloying treatment is performed at a temperature exceeding 600 ° C., untransformed austenite is transformed into pearlite, and a desired volume ratio of retained austenite cannot be secured, and El may be lowered. Therefore, when the galvanizing alloying treatment is performed, it is preferable to perform the galvanizing alloying treatment in a temperature range of 470 ° C. or more and 600 ° C. or less. Moreover, you may perform an electrogalvanization process. In addition, the plating adhesion amount is preferably 20 to 80 g / m ² per side (double-sided plating), and the alloyed hot-dip galvanized steel sheet (GA) is subjected to alloying treatment so that the Fe concentration in the plating layer is 7 to 15 mass. % Is preferable.

The reduction ratio of the skin pass rolling after the heat treatment is preferably in the range of 0.1% to 2.0%. If it is less than 0.1%, the effect is small and control is difficult, so this is the lower limit of the good range. Moreover, since productivity will fall remarkably when it exceeds 2.0%, this is made the upper limit of a favorable range.
Skin pass rolling may be performed online or offline. Further, a skin pass having a desired reduction rate may be performed at once, or may be performed in several steps. Other manufacturing method conditions are not particularly limited, but from the viewpoint of productivity, a series of treatments such as annealing, hot dip galvanization, alloying treatment of galvanization, etc. are performed by CGL (Continuous Galvanizing) which is a hot dip galvanizing line. Line). After hot dip galvanization, wiping is possible to adjust the amount of plating. In addition, conditions, such as plating other than the above-mentioned conditions, can depend on the conventional method of hot dip galvanization.

Example 1
Steel having the component composition shown in Table 1 and the balance being Fe and inevitable impurities was melted in a converter and made into a slab by a continuous casting method. The obtained slab was heated under the conditions shown in Table 2 and hot-rolled, and then pickled. Nos. 1 to 11, 13 to 23, 25, 27, 29, 30, 32 to 37, 39, and 41 were subjected to hot-rolled sheet heat treatment. Nos. 29, 30, 32 to 37, 39, and 41 were subjected to pickling treatment after the heat treatment of the plate.
Next, after cold rolling under the conditions shown in Table 2, annealing was performed twice under the conditions shown in Table 2 to obtain a high-strength cold-rolled steel sheet (CR).
Furthermore, some high-strength cold-rolled steel sheets (CR) were galvanized to obtain hot-dip galvanized steel sheets (GI), galvannealed steel sheets (GA), electrogalvanized steel sheets (EG), and the like. The hot dip galvanizing bath uses a zinc bath containing Al: 0.14% by mass or 0.19% by mass in GI, and uses a zinc bath containing Al: 0.14% by mass in GA. Was 470 ° C. Coating weight, the GI, and per side 72 g / m ² or 45 g / m ² (two-sided plating), also in GA, and per one surface 45 g / m ² (two-sided plating). Moreover, GA made Fe density | concentration in a plating layer 9 mass% or more and 12 mass% or less. Furthermore, the amount of EG plating adhered was 50 g / m ² per side (double-sided plating).

In addition, T1 temperature (degreeC) was calculated | required using the following formula | equation.
T1 temperature (° C.) = 720 + 29 × [% Si] -21 × [% Mn] + 17 × [% Cr]
The T2 temperature (° C) is
T2 temperature (° C.) = 946−203 × [% C] ^1/2 + 45 × [% Si] −30 × [% Mn] + 150 × [% Al] −20 × [% Cu] + 11 × [% Cr] +350 × [% Ti] + 104 × [% V]
Can be calculated. Here, [% X] is the mass% of the component element X of the steel sheet, and zero for the component elements not contained.
T1 means Ac ₁ point and T2 means Ac ₃ point.

The high-strength cold-rolled steel sheet (CR), hot-dip galvanized steel sheet (GI), alloyed hot-dip galvanized steel sheet (GA) and electrogalvanized steel sheet (EG) obtained as described above were used as test steels. Characteristics were evaluated. The mechanical properties were evaluated by performing a tensile test and Young's modulus measurement as follows. The results are shown in Table 3. Table 3 also shows the thickness of each steel plate as the test steel.

The tensile test is based on JIS Z 2241 (2011) using a JIS No. 5 test piece obtained by taking a sample so that the length of the tensile test piece is perpendicular to the rolling direction of the steel sheet (C direction). And TS (tensile strength) and El (total elongation) were measured. In addition, in this invention, it is a case where the value of TSxEl is 15000 MPa *% or more that it is excellent in ductility, ie El.

Young's modulus measurement is 10 mm × 50 mm from three directions, ie, the rolling direction of the steel sheet (L direction), the 45 ° direction (D direction) with respect to the rolling direction of the steel sheet, and the direction perpendicular to the rolling direction of the steel sheet (C direction). The test piece was cut out and the Young's modulus was measured using a transverse vibration type resonance frequency measuring device according to the American Society to Testing Materials standard (C1259). In the present invention, the rigidity, that is, the Young's modulus is excellent when the Young's modulus in the 45 ° direction with respect to the rolling direction and the rolling direction is 205 GPa or more and the Young's modulus in the direction perpendicular to the rolling direction is 220 GPa or more. is there.

The average r-value measurement was taken from three directions, ie, the rolling direction (L direction) of the steel sheet, the 45 ° direction (D direction) with respect to the rolling direction of the steel sheet, and the direction perpendicular to the rolling direction of the steel sheet (C direction). Using the JIS No. 5 test piece, each plastic strain ratio r _L , r _D , r _C was determined according to JIS Z 2254 (2008), and the average r value was calculated by the following formula.
Average r value = (r _L + 2r _D + r _C ) / 4
Note that the present invention is excellent in deep drawability when the average r value, which is an index of deep drawability, is 0.95 or more regardless of the strength of the steel sheet.

Further, according to the above-described method, the area ratio of ferrite, bainite and martensite, the volume ratio of retained austenite, and the inverse strength ratio of γ-fiber to α-fiber of ferrite at the 1/4 thickness position of the steel sheet were obtained. The results are also shown in Table 3.

As shown in Table 3, in the inventive examples, TS is 780 MPa or more, excellent in ductility, has a balance between high strength and ductility, and is excellent in rigidity and deep drawability. On the other hand, in the comparative example, one or more of strength, ductility, balance between strength and ductility, rigidity, and deep drawability was inferior.
As mentioned above, although embodiment of this invention was described, this invention is not limited by the description which makes a part of indication of this invention by this embodiment. That is, all other embodiments, examples, operation techniques, and the like made by those skilled in the art based on the present embodiment are all included in the technical scope of the present invention. For example, in the series of heat treatments in the above-described manufacturing method, as long as the heat history condition is satisfied, the equipment for performing the heat treatment on the steel sheet is not particularly limited.

According to the present invention, it is possible to produce a high-strength steel sheet having a TS of 780 MPa or more and excellent not only in ductility but also in rigidity and also in deep drawability. Therefore, by applying the high-strength steel plate obtained by the present invention to, for example, an automobile structural member, fuel efficiency can be improved by reducing the weight of the vehicle body, and the industrial utility value is extremely large.

Claims (4)

1. Ingredient composition is mass%,
   C: 0.08% to 0.35%,
   Si: 0.50% or more and 2.50% or less,
   Mn: 1.50% or more and 3.00% or less,
   P: 0.001% to 0.100%,
   S: 0.0001% or more and 0.0200% or less and N: 0.0005% or more and 0.0100% or less,
   Ti: 0.001% or more and 0.200% or less and Nb: 0.001% or more and 0.200% or less containing one or two selected from the balance, the balance consists of Fe and inevitable impurities,
   Steel structure is area ratio,
   Ferrite is 20% or more, the average grain size of the ferrite is 10 μm or more and 20 μm or less, and bainite is 5% or more,
   Martensite is 5% or more,
   The total area ratio of bainite and martensite is 15% or more,
   In volume ratio, residual austenite is 5% or more,
   Further, the ferrite texture has a microstructure in which the inverse strength ratio of γ-fiber to α-fiber is more than 3.0.
   High strength steel plate.
2. In the high-strength steel sheet according to claim 1, further in mass%,
   Al: 0.01% or more and 1.00% or less,
   V: 0.005% or more and 0.100% or less,
   B: 0.0001% to 0.0050%,
   Cr: 0.05% or more and 1.00% or less,
   Cu: 0.05% or more and 1.00% or less,
   Sb: 0.0020% or more and 0.2000% or less,
   Sn: 0.0020% or more and 0.2000% or less,
   Ta: 0.0010% or more and 0.1000% or less,
   Ca: 0.0003% or more and 0.0050% or less,
   A high-strength steel sheet containing at least one element selected from Mg: 0.0003% to 0.0050% and REM: 0.0003% to 0.0050%.
3. It is a manufacturing method of the high strength steel plate according to claim 1 or 2,
   The steel slab having the component composition according to claim 1 or 2 is heated to 1100 ° C or higher and 1300 ° C or lower, hot rolled at a finish rolling exit temperature: 800 ° C or higher and 1000 ° C or lower, and a coiling temperature: 300 ° C. Winding at 800 ° C. or lower, after pickling treatment, or after holding at 900 ° C. or higher and 36000 s or lower in the temperature range of 450 ° C. or higher and 800 ° C. or lower, then cold rolling at a rolling reduction of 40% or higher,
   Subsequently, the obtained cold-rolled sheet is heated to a temperature range of 450 ° C. or more and T1 temperature or less, and subjected to a first annealing treatment for holding at this temperature range for 300 s or more,
   Next, after reheating to a temperature range of T1 temperature or more and T2 temperature or less and performing the second annealing treatment, cooling is performed at 300 ° C or more and 500 ° C or less at an average cooling rate of at least 550 ° C at 5 ° C / s or more. A method for producing a high-strength steel sheet, which is cooled to a stop temperature range and held for 10 s or longer in the cooling stop temperature range.
   T1 temperature (° C.) = 720 + 29 × [% Si] -21 × [% Mn] + 17 × [% Cr]
   T2 temperature (° C.) = 946−203 × [% C] ^1/2 + 45 × [% Si] −30 × [% Mn] + 150 × [% Al] −20 × [% Cu] + 11 × [% Cr] +350 × [% Ti] + 104 × [% V]
   [% X] is the mass% of the component element X of the steel sheet, and zero for the component elements not contained.
4. A high-strength galvanized steel sheet having a galvanized layer on the surface of the high-strength steel sheet according to claim 1 or 2.