JP2019011508A - Low-yield-ratio high strength steel sheet and method for producing the same - Google Patents

Abstract

To provide a low-yield-ratio high strength steel sheet having not more than 0.70 of a yield ratio and 590 MPa or more of a tensile strength, and having excellent moldability and chemical conversion treatability.SOLUTION: A high-yield-ratio high strength steel sheet contains, in mass%, C: more than 0.050% to 0.200% or less, Si: 0.01% or more to less than 0.50%, Mn: more than 1.80% to less than 2.60%, P: 0.100% or less, S: 0.0200% or less, Al: 2.000% or less, N: 0.0100% or less, and Ti: 0.005% or more to 0.100% or less and/or Nb: 0.005% or more and 0.100% or less with the balance being Fe and inevitable impurities, wherein, in area ratios, ferrite is 40.0-90.0%, martensite is 5.0-30.0%, and perlite is 2.0-30.0%, with Mn content in steel/Mn content in ferrite≥1.10 or more, Mn content in martensite/Mn content in steel≥1.80 or more, wherein a measure of Mn content is mass%.SELECTED DRAWING: None

Description

The present invention relates to a low yield ratio type high-strength steel sheet and a method for producing the same.
More specifically, the present invention relates to a member or component used in the industrial field such as an automobile, an electric machine, a pipe, and a container, particularly a low yield ratio type high-strength steel plate suitable for an automobile member and a method for manufacturing the same.

In recent years, from the viewpoint of conservation of the global environment, there has been a strong demand for improvement in fuel consumption for the purpose of reducing automobile CO ₂ emissions. For this reason, efforts have been made to reduce the thickness of the vehicle body by increasing the strength of the vehicle body material and to reduce the weight of the vehicle body itself. For example, cold-rolled steel sheets used for parts manufactured by press forming are often used as high-strength steel sheets having a tensile strength (TS) of 590 MPa or more (for example, Patent Documents 1 to 3). See).

  A steel plate having a TS of 590 MPa or more may be modularized by assembling by spot welding, arc welding, or the like after press working in an automobile manufacturing process. In this case, high dimensional accuracy is required during assembly. In such a steel plate, it is necessary to make it difficult for spring back to occur after processing, so it is important that the yield ratio (YR) is low before processing.

JP 2003-342680 A JP 2011-2119855 A Japanese Patent Laid-Open No. 2006-141857

In general, increasing the strength of a steel sheet causes a decrease in formability (ductility, deep drawability, hole expansibility). For this reason, as the strength of the steel sheet increases, cracks during forming may become a problem.
In particular, in a high-strength steel sheet having a TS of 590 MPa or more, it is difficult to improve ductility, deep drawability, and hole expandability in a superimposed manner. Therefore, development of a steel sheet that has both high strength and excellent formability (ductility, deep drawability, hole expansibility) is desired.

  In addition, cold-rolled steel sheets used for automobiles and the like are used after being coated, but are subjected to chemical conversion treatment as pretreatment for ensuring corrosion resistance after painting. A chemical conversion film is formed by chemical conversion treatment. At this time, there may be a micro region (scaling) in which no chemical conversion film is formed (that is, chemical conversion processability is insufficient), and good chemical conversion processability is required. It is done.

In addition, as annealing at the time of obtaining a high-strength steel plate having a TS of 590 MPa or more, continuous annealing by CAL (Continuous Annealing Line) is mainstream from the viewpoint of manufacturing lead time.
However, continuous annealing by CAL is limited in production base due to large-scale equipment, and the plate thickness and width that can be passed through are limited as the strength of the steel plate increases, and as the strength of the steel plate increases Coil joint welding may be difficult due to the large amount of alloying elements added.
For this reason, it may be required to obtain a high-strength steel sheet having a TS of 590 MPa or more by batch annealing using BAF (Batch Annealing Furnace).

When the present inventors manufactured a steel plate through batch annealing by BAF using a steel slab having a component composition (C: 0.05% by mass or less) described in Patent Document 1, a desired amount of martensite is obtained. May not be secured, and TS of 590 MPa or more may not be obtained.
In addition, the present inventors have a component composition described in Patent Document 2 (Si: more than 0.5% by mass) and a component composition described in Patent Document 3 (Si: 0.50% by mass or more). When a steel plate was manufactured through batch annealing with BAF using a slab, chemical conversion treatment was sometimes insufficient.

The present invention has been made in view of the above points, and has a low yield ratio (YR) of 0.70 or less and a low yield ratio type high strength steel sheet having a tensile strength (TS) of 590 MPa or more. An object of the present invention is to provide a low-yield ratio type high-strength steel sheet that is excellent in formability (ductility, deep drawability, hole expansibility) and chemical conversion treatment.
Furthermore, this invention also aims at providing the manufacturing method of the low yield ratio type high strength steel plate which obtains the said low yield ratio type high strength steel plate by batch annealing.

  As a result of intensive studies by the present inventors, it has been found that the above object can be achieved by adopting the following constitution, and the present invention has been completed.

That is, the present invention provides the following [1] to [5].
[1] By mass%, C: more than 0.050% and 0.200% or less, Si: 0.01% or more and less than 0.50%, Mn: more than 1.80% and less than 2.60%, P: 0.00. 100% or less, S: 0.0200% or less, Al: 2.000% or less, and N: 0.0100% or less, and Ti: 0.005% or more and 0.100% or less, and Nb: containing at least one selected from the group consisting of 0.005% or more and 0.100% or less, with the component composition of the balance consisting of Fe and inevitable impurities and the area ratio, ferrite is 40.0% or more and 90 0.0% or less, martensite 5.0% or more and 30.0% or less, and pearlite 2.0% or more and 30.0% or less. The value divided by the amount of Mn is 1.10 or more. Mn amount is the value obtained by dividing the Mn content in the steel is 1.80 or more, the Mn amount of the unit is mass%, the low yield ratio high-strength steel sheets in the.
[2] The low yield ratio type high-strength steel sheet according to the above [1], wherein the component composition further contains at least one selected from the following group A to group E by mass%.
[Group A] Ni: 0.01% or more and 1.00% or less, Cu: 0.005% or more and 1.000% or less, Cr: 0.01% or more and 1.00% or less, and Mo: 0.005 %: At least one selected from the group consisting of 0.5% to 0.500% [Group B] V: 0.005% to 0.100% and W: 0.005% to 0.100% [Group C] Ca: 0.0005% or more and 0.0050% or less, Mg: 0.0005% or more and 0.0050% or less, and REM: 0.0005% or more and 0.0050% or less At least one selected from the group consisting of [Group D] Sb: 0.002% to 0.200%, Sn: 0.002% to 0.200%, and Ta: 0.001% to 0.000. Selected from the group consisting of 010% or less At least one [E group] B: 0.0001% or more 0.0050% or less [3] further satisfies the following formulas, the above-mentioned [1] or [2] Low yield ratio high-strength steel sheet according to.
P (222) / {P (200) + P (220)} ≧ 2.0
In the above formula, P (222), P (200), and P (220) are respectively the (222) plane, the (200) plane, and the (220) plane parallel to the plate plane at the ¼ plate thickness position. Represents the diffraction X-ray integral intensity ratio.
[4] The low yield ratio type high-strength steel sheet according to any one of [1] to [3], wherein a volume ratio of retained austenite in the steel structure is 5.0% or less.
[5] A hot-rolled steel sheet is obtained by subjecting the steel slab having the composition described in [1] or [2] above to hot rolling at a finish rolling exit temperature of 800 ° C. or higher and 1000 ° C. or lower, The hot rolled steel sheet is cooled at an average cooling rate of 20 ° C./second to 120 ° C./second in the temperature range from the finish rolling exit temperature to 650 ° C., and the temperature range from 650 ° C. to the following winding temperature is set. The steel sheet is cooled at an average cooling rate of 5 ° C./second or more and 40 ° C./second or less, the hot-rolled steel sheet after the cooling is wound at a winding temperature of 400 ° C. or more and 600 ° C. or less, and the wound hot-rolled steel sheet is wound. The steel sheet is pickled, and the hot-rolled steel sheet subjected to the pickling is cold-rolled at a rolling reduction of 30% or more and 85% or less to obtain a cold-rolled steel sheet, and the cold-rolled steel sheet Is subjected to batch annealing satisfying the following temperature histories 1 and 2, and By subjecting the cold-rolled steel sheet that has been annealed to temper rolling at an elongation of 1.1% or less, the low yield ratio type high-strength steel sheet according to any one of the above [1] to [4]. A method for producing a low yield ratio type high strength steel sheet.
Temperature history 1: The position of 5 mm along the radial direction from the outer peripheral surface to the inner peripheral surface of the coil of the cold-rolled steel sheet is heated to 620 ° C. or higher and 760 ° C. or lower for 2.0 hours or more and 72.0 or more. The temperature range of 400 ° C. or more and 550 ° C. or less is kept at an average cooling rate of 30 ° C./hour or more and 200 ° C./hour or less.
Temperature history 2: After the temperature was raised from the inner peripheral surface to the outer peripheral surface of the coil of the cold-rolled steel sheet, the position where 1/3 of the coil thickness was entered was 2.0 ° C. to 600 ° C. to 740 ° C. The temperature range of 400 ° C. to 550 ° C. is maintained at an average cooling rate of 5 ° C./hour to 100 ° C./hour.

According to the present invention, a low yield ratio type high strength steel sheet having a low yield ratio (YR) of 0.70 or less and a tensile strength (TS) of 590 MPa or more, which has formability (ductility, depth). It is possible to provide a low yield ratio type high-strength steel sheet that is excellent in drawability, hole expansibility) and chemical conversion treatment.
Furthermore, according to this invention, the manufacturing method of the low yield ratio type high strength steel plate which obtains the said low yield ratio type high strength steel plate by batch annealing can be provided.

It is a perspective view which shows the coil which notched a part.

In the present specification, a numerical range expressed using “to” means a range including numerical values described before and after “to” as a lower limit value and an upper limit value.
Hereinafter, the low yield ratio type high strength steel plate is also simply referred to as “high strength steel plate” or “steel plate”.

[Knowledge obtained by the present inventors]
The present inventors have a low yield ratio (YR) of 0.70 or less, a tensile strength (TS) of 590 MPa or more, and formability (ductility, deep drawability, hole expansibility). In addition, in order to develop a low-yield ratio type high-strength steel plate that can be manufactured by batch annealing with BAF, which has excellent chemical conversion properties, it has been intensively studied. As a result, the following knowledge was obtained.

(1) The following points are important for obtaining TS of 590 MPa or more in the batch annealing by BAF.
That is, Mn is contained in a range exceeding 1.80% by mass, and batch annealing by BAF is performed in a two-phase region of ferrite and austenite. By maintaining for a long time in the two-phase region, not only the interstitial element C but also the substitutional element Mn is concentrated from ferrite into austenite. Thereby, even if it is a very slow cooling rate after batch annealing by BAF, a desired amount of martensite can be secured.

(2) The following points are important for obtaining excellent formability (ductility, deep drawability, hole expansibility) in production by batch annealing with BAF.
In order to obtain excellent ductility, high work hardening ability expressed from the DP (Dual Phase) structure of ferrite and martensite is obtained. Therefore, the amount of ferrite and martensite is adjusted appropriately. Furthermore, the amount of Mn in the ferrite is reduced, the ferrite is cleaned, and a ferrite rich in ductility is obtained.
In order to obtain excellent deep drawability, an appropriate amount of ferrite is secured and a {111} recrystallized texture of the ferrite is developed. Furthermore, the amount of Mn in the ferrite is reduced and the ferrite is cleaned. Both are realized by long-time heat treatment in batch annealing by BAF. By similar structure control and manufacturing method, the Young's modulus in the direction perpendicular to the rolling direction of the steel sheet (C direction) is also improved.
In order to obtain excellent hole expandability, the amount of boundary between the void starting point and the soft phase ferrite and the hard phase martensite, which are crack propagation paths, is adjusted. Therefore, the amount of martensite is controlled appropriately. Furthermore, by actively utilizing pearlite, which is an intermediate hardness phase between ferrite, which is a soft phase, and martensite, which is a hard phase, the boundary amount between ferrite and martensite is reduced while securing strength. Any of them can be adjusted by the soaking time and soaking temperature and the average cooling rate during the subsequent cooling in batch annealing by BAF.

(3) The following points are important in order to have good chemical conversion properties in the production by batch annealing with BAF.
Si amount is made less than 0.50%. Furthermore, when the soaking temperature in batch annealing is too high, concentration of Mn on the surface of the steel sheet during annealing is promoted, and chemical conversion treatment performance is deteriorated. For this reason, an appropriate upper limit temperature is set for the soaking temperature of batch annealing.

(4) In order to obtain a low YR of 0.70 or less, the following points are important.
After batch annealing by BAF, tempering (skin pass) rolling with an elongation of 1.1% or less is performed. A steel sheet produced by batch annealing with BAF includes many boundaries between ferrite and martensite, and thus a low YR is obtained. Thereafter, YR can be controlled low by introducing a small amount of strain in temper rolling with an elongation of 1.1% or less. This can be adjusted by off-line temper rolling at room temperature after batch annealing with BAF.

  (5) Furthermore, in order to build the structure as described above, it is important to adjust the component composition to a predetermined range and appropriately control the manufacturing conditions, particularly the conditions for batch annealing by BAF.

The present invention was completed after further studies based on the above findings.
In the following, first, the low yield ratio type high strength steel sheet of the present invention will be described, and then the manufacturing method of the low yield ratio type high strength steel sheet of the present invention will be described.

[Low yield ratio type high strength steel sheet]
The low yield ratio type high-strength steel sheet of the present invention (hereinafter also simply referred to as “the high-strength steel sheet of the present invention” or “the steel sheet of the present invention”) is in mass% and C: more than 0.050% and 0.200% or less. , Si: 0.01% or more and less than 0.50%, Mn: more than 1.80% and less than 2.60%, P: 0.100% or less, S: 0.0200% or less, Al: 2.000% or less N: 0.0100% or less, and at least selected from the group consisting of Ti: 0.005% to 0.100% and Nb: 0.005% to 0.100% 1 type is contained, the remainder is composed of Fe and inevitable impurities, and the area ratio, ferrite is 40.0% or more and 90.0% or less, martensite is 5.0% or more and 30.0% or less, A steel structure whose pearlite is 2.0% or more and 30.0% or less; To. Further, in the low yield ratio type high strength steel sheet of the present invention, the value obtained by dividing the Mn content in the steel by the Mn content in the ferrite is 1.10 or more, and the Mn content in the martensite is changed to Mn in the steel. The value divided by the amount is 1.80 or more. The unit of the amount of Mn is mass%.

  The low yield ratio type high strength steel sheet of the present invention has a yield ratio (YR) of 0.70 or less and a tensile strength (TS) of 590 MPa or more, and further has formability (ductility, deep drawability, Excellent hole expandability) and chemical conversion treatment.

In the present invention, the “low yield ratio type” means that the yield ratio (YR) is 0.70 or less. YR is a value obtained by dividing the yield stress (YS) by the tensile strength (TS). YR is preferably 0.50 or more and 0.68 or less.
In the present invention, “high strength” means that the tensile strength (TS) is 590 MPa or more.

  The high-strength steel sheet of the present invention can be applied to, for example, automobile members.

  The plate thickness of the high-strength steel plate of the present invention is not particularly limited, and is, for example, 0.5 mm or more and 4.0 mm or less.

<Ingredient composition>
Below, the component composition of the high strength steel plate of this invention is demonstrated first. Unless otherwise specified, “%” in the component composition means “mass%”.

<< C: more than 0.050% and 0.200% or less >>
C is an element necessary for generating martensite and increasing the strength.
If the amount of C is 0.050% or less, the amount of ferrite increases, it is difficult to ensure the desired amount of martensite, and the desired strength cannot be obtained.
On the other hand, if the amount of C exceeds 0.200%, the amount of hard martensite becomes excessive, and the hole expandability and the like deteriorate. Hardening of the welded part and the heat-affected zone becomes remarkable, the mechanical properties of the welded part are lowered, and spot weldability, arc weldability, and the like may be deteriorated.
For this reason, the amount of C is more than 0.050% and 0.200% or less, and preferably 0.065% or more and 0.150% or less.

<< Si: 0.01% or more and less than 0.50% >>
Si is an element effective for ensuring good ductility because it improves the work hardening ability of ferrite. If the amount of Si is less than 0.01%, the effect of addition becomes poor.
On the other hand, excessive addition of 0.50% or more of Si results in a decrease in the chemical conversion property of the cold-rolled steel sheet or a temper color. The reason why the chemical conversion treatment performance is lowered is that Si is very easy to oxidize and forms a film of Si oxide (SiO ₂ ) on the surface of the steel sheet. This Si oxide inhibits the formation reaction of the chemical conversion film during the chemical conversion treatment. However, it is presumed that a micro region (scaling) where a chemical conversion film is not generated is generated. The temper color is a surface defect caused by a surface oxide film generated by oxidation of a component in the steel (particularly, Si) that is concentrated on the surface of the steel sheet during annealing. Furthermore, it causes deterioration of the surface properties due to the occurrence of red scale and the like.
For this reason, Si amount is 0.01% or more and less than 0.50%, and 0.10% or more and 0.40% or less are preferable.

<< Mn: more than 1.80% and less than 2.60% >>
Mn is an extremely important element in the present invention. Mn is an element that stabilizes austenite. When holding for a long time such as batch annealing by BAF in the two-phase region of ferrite and austenite, Mn is concentrated in the austenite, and even if the cooling rate is very slow, almost no ferrite transformation or bainite transformation occurs. The amount of martensite can be secured. Such an effect is recognized when the amount of Mn in steel exceeds 1.80%. Thereby, the stable manufacture whose tensile strength (TS) is 590 MPa or more is possible by batch annealing.
On the other hand, when Mn is added excessively in an amount of 2.60% or more, martensite with an area ratio exceeding 30.0% is generated, and desired ductility cannot be obtained.
For this reason, the amount of Mn is more than 1.80% and less than 2.60%, and preferably 2.00% or more and less than 2.60%.

<< P: 0.100% or less >>
P is an element that has a solid solution strengthening action and can be added according to a desired strength. P is an element that promotes ferrite transformation and is also effective for forming a composite structure of a steel sheet. In order to obtain such an effect, the amount of P is, for example, 0.001% or more.
However, if the P content exceeds 0.100%, the spot weldability is significantly deteriorated. For this reason, the amount of P is 0.100% or less, and preferably 0.040% or less.

<< S: 0.0200% or less >>
S segregates at the grain boundaries and embrittles the steel during hot working, and also exists as a sulfide and lowers the local deformability of the steel sheet. When the amount of S exceeds 0.0200%, the spot weldability is significantly deteriorated. For this reason, S amount is 0.0200% or less, 0.0100% or less is preferable and 0.0050% or less is more preferable.
The lower limit of the amount of S is not particularly limited, but the amount of S is, for example, 0.0001% or more due to restrictions on production technology.

<< Al: 2.000% or less >>
Al is an element effective in expanding the two-phase region of ferrite and austenite and reducing temperature dependency during annealing, that is, material stability. Al acts as a deoxidizer and is also an effective element for the cleanliness of steel. For this reason, Al amount is 0.005% or more, for example, and 0.010% or more is preferable.
However, the addition of a large amount of Al increases the risk of cracking of steel slabs during continuous casting, and decreases manufacturability. For this reason, the amount of Al is 2.000% or less, and preferably 1.500% or less.

<< N: 0.0100% or less >>
N is an element that degrades the aging resistance of steel. In particular, when the N content exceeds 0.0100%, the deterioration of aging resistance becomes significant. For this reason, N amount is 0.0100% or less, and 0.0070% or less is preferable.
The N content is preferably as small as possible, but is, for example, 0.0005% or more and preferably 0.0010% or more because of restrictions on production technology.

  The component composition of the high-strength steel sheet of the present invention further includes at least one selected from the group consisting of Ti: 0.005% to 0.100% and Nb: 0.005% to 0.100%. contains.

<< Ti: 0.005% or more and 0.100% or less >>
Ti is an extremely important element in the present invention. Ti forms precipitates with C, S, and N, and produces a ferrite having an orientation that is advantageous for improving deep drawability and Young's modulus (rigidity) during annealing. Furthermore, Ti suppresses the coarsening of recrystallized grains and contributes effectively to improving the strength. 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 addition, ductility at high temperatures is improved, and castability in continuous casting is improved. These effects are obtained when the Ti content is 0.005% or more.
On the other hand, when the Ti content exceeds 0.100%, the carbonitride content is remarkably increased and the ductility is lowered.
For this reason, Ti amount is 0.005% or more and 0.100% or less, and preferably 0.020% or more and 0.090% or less.

<< Nb: 0.005% or more and 0.100% or less >>
Nb is an extremely important element in the present invention. Nb forms fine precipitates at the time of hot rolling or annealing, and generates a ferrite having an orientation that is advantageous for improving deep drawability and Young's modulus (rigidity) during annealing. Furthermore, Nb suppresses the coarsening of recrystallized grains and contributes effectively to improving the strength. These effects are obtained when the Nb amount is 0.005% or more.
On the other hand, when the Nb content exceeds 0.100%, the carbonitride content is remarkably increased and the ductility is lowered. It also becomes a factor of cost increase.
For this reason, the amount of Nb is 0.005% or more and 0.100% or less, and preferably 0.010% or more and 0.080% or less.

  The component composition of the high-strength steel sheet of the present invention can further contain at least one selected from the following groups A to E in mass%.

<< [Group A] Ni: 0.01% to 1.00%, Cu: 0.005% to 1.000%, Cr: 0.01% to 1.00%, and Mo: 0.0. At least one selected from the group consisting of 005% and 0.500% >>
Ni is an element that stabilizes retained austenite and is effective in ensuring good ductility. Further, Ni is an element that increases the strength of steel by solid solution strengthening. From the viewpoint of obtaining this addition effect, the Ni content is preferably 0.01% or more. On the other hand, if the Ni content exceeds 1.00%, the area ratio of hard martensite may be excessive. It also becomes a factor of cost increase. For this reason, when adding Ni, the amount of Ni is preferably 0.01% or more and 1.00% or less.
Cu is an element effective for increasing the strength of steel. From the viewpoint of obtaining this addition effect, the amount of Cu is preferably 0.005% or more. On the other hand, if the amount of Cu exceeds 1.000%, the amount of hard martensite may be excessive. For this reason, when adding Cu, the amount of Cu is preferably 0.005% or more and 1.000% or less.
Cr and Mo are elements that increase the strength of the steel and contribute to improving the hardenability. From the viewpoint of obtaining this addition effect, the Cr content is preferably 0.01% or more, and the Mo content is preferably 0.005% or more. On the other hand, if these elements are added excessively, the amount of hard martensite may become excessive. It also becomes a factor of cost increase. For this reason, when adding Cr, Cr amount is preferably 0.01% or more and 1.00% or less, and Mo amount is preferably 0.005% or more and 0.500% or less.

<< [Group B] V: 0.005% or more and 0.100% or less; and W: at least one selected from the group consisting of 0.005% or more and 0.100% or less >>
V and W are elements that can be added as necessary because they are effective for strengthening steel by precipitation strengthening. From the viewpoint of obtaining this addition effect, the V amount is preferably 0.005% or more, and the W amount is preferably 0.005% or more. On the other hand, when these elements are added excessively, the amount of carbonitride is remarkably increased, and the ductility may be lowered. It also becomes a factor of cost increase.
Therefore, when adding at least one selected from the group consisting of V and W, the V amount is preferably 0.005% or more and 0.100% or less, and the W amount is 0.005% or more and 0.100% or less. preferable.

<< [C Group] Ca: 0.0005% or more and 0.0050% or less; Mg: 0.0005% or more and 0.0050% or less; and REM: 0.0005% or more and 0.0050% or less. At least one kind
Ca, Mg and REM are effective elements for spheroidizing the shape of the sulfide and improving the adverse effect of the sulfide on the hole expandability. From the viewpoint of obtaining this effect of addition, the amount of each element is preferably 0.0005% or more. On the other hand, if each of Ca, Mg and REM is added excessively, inclusions and the like are increased, which may cause surface and internal defects.
For this reason, when adding at least 1 sort (s) chosen from the group which consists of Ca, Mg, and REM, the amount is preferably 0.0005% or more and 0.0050% or less, respectively.

<< [Group D] Sb: 0.002% to 0.200%, Sn: 0.002% to 0.200%, and Ta: 0.001% to 0.010% At least one kind
Sn and Sb are elements that can be added as necessary from the viewpoint of suppressing decarburization in the thickness region of about several tens of μm of the steel sheet surface layer caused by nitriding or oxidizing the steel sheet surface. By suppressing such nitriding and oxidation, the amount of martensite on the steel sheet surface can be prevented from decreasing, so Sn and Sb are effective elements for ensuring strength and material stability. On the other hand, when Sn and Sb are added excessively, the toughness may be lowered. For this reason, when adding at least 1 sort (s) chosen from the group which consists of Sn and Sb, the amount is respectively 0.002% or more and 0.200% or less.
Ta, like Ti and Nb, generates alloy carbide and alloy carbonitride and contributes to high strength. In addition, Ta partially dissolves in Nb carbides and Nb carbonitrides to form composite precipitates such as (Nb, Ta) (C, N), thereby suppressing the coarsening of precipitates and strengthening precipitation. It is thought that there is an effect that stabilizes the contribution to the strength improvement by. From the viewpoint of obtaining the precipitate stabilization effect described above, the Ta amount is preferably 0.001% or more. On the other hand, when Ta is added excessively, the effect of addition is saturated and the cost also increases. For this reason, when adding Ta, the amount of Ta is preferably 0.001% or more and 0.010% or less.

<< [Group E] B: 0.0001% or more and 0.0050% or less >>
B has an action of suppressing the formation and growth of ferrite from the austenite grain boundary, and can be added as necessary because it can flexibly control the structure. For example, there is an effect of suppressing ferrite transformation and pearlite transformation during primary cooling after hot rolling. Furthermore, it suppresses ferrite transformation and bainite transformation during cooling after holding in batch annealing. From the viewpoint of obtaining such an effect of addition, the B content is preferably 0.0001% or more.
On the other hand, if the amount of B exceeds 0.0050%, formability (ductility, deep drawability, hole expansibility) may be deteriorated.
For this reason, when adding B, B amount is preferably 0.0001% or more and 0.0050% or less, and more preferably 0.0005% or more and 0.0030% or less.

《Remainder》
In the above component composition, the balance other than the above components is Fe and inevitable impurities.

<Steel structure>
Next, the steel structure (microstructure) of the high-strength steel sheet of the present invention will be described.

<< Area ratio of ferrite: 40.0% or more and 90.0% or less >>
In the high-strength steel sheet of the present invention, the area ratio of ferrite is set to 40.0% or more in order to ensure desired ductility and deep drawability. When the area ratio of ferrite is 40.0% or more, the amount of ferrite that is a soft phase is sufficient, and excellent ductility can be secured. By developing a {111} recrystallized texture of ferrite having an area ratio of 40.0% or more, a desired deep drawability can be secured.
On the other hand, in order to secure TS of 590 MPa or more, the area ratio of soft ferrite is set to 90.0% or less.
For this reason, the area ratio of ferrite is 40.0% or more and 90.0% or less, and preferably 50.0% or more and 85.0% or less.

  The {111} texture means that the <111> direction of the crystal is oriented in the direction perpendicular to the steel plate surface. From the crystallographic and Bragg reflection conditions, in the case of α-Fe having a body-centered cubic structure, diffraction on the {111} plane does not occur on the (111) plane but occurs on the (222) plane. As the intensity ratio, the value (P (222)) of the (222) plane is used. The (222) plane is substantially the same as the <111> direction because the [222] direction is oriented in the direction perpendicular to the steel plate face. Therefore, a high intensity ratio of the (222) plane corresponds to the development of the {111} texture. For the same reason, the value of the (200) plane (P (200)) is used for the {100} plane.

<< Martensite area ratio: 5.0% or more and 30.0% or less >>
In the high-strength steel sheet of the present invention, in order to achieve a TS of 590 MPa or more, the martensite area ratio is set to 5.0% or more.
On the other hand, in order to ensure desired ductility and hole expandability, the martensite area ratio is set to 30.0% or less. When the area ratio of martensite is 30.0% or less, the amount of a phase softer than martensite other than martensite increases, and thus excellent ductility can be secured. When the area ratio of martensite is 30.0% or less, since the boundary amount between ferrite, which is a soft phase, and martensite, which is a hard phase, is small, excellent hole expandability can be secured.
For this reason, the area ratio of martensite is 5.0% or more and 30.0% or less, and preferably 8.0% or more and 25.0% or less.

<< Perlite area ratio: 2.0% or more and 30.0% or less >>
In the high-strength steel sheet of the present invention, the area ratio of pearlite is set to 2.0% or more in order to ensure a desired hole expanding property. When the area ratio of pearlite is 2.0% or more, by using pearlite, which is an intermediate hardness phase between ferrite, which is a soft phase, and martensite, which is a hard phase, while maintaining strength, ferrite and martensite As a result, it is possible to reduce the amount of boundary between the holes, so that excellent hole expandability can be secured.
On the other hand, in order to ensure the desired ductility, the area ratio of pearlite is set to 30.0% or less. If the area ratio of pearlite is 30.0% or less, excellent ductility can be ensured by the high work hardening ability expressed from the remaining ferrite and martensite DP (Dual Phase) structures.
For this reason, the area ratio of pearlite is 2.0% or more and 30.0% or less, and preferably 4.0% or more and 25.0% or less.

The area ratio of ferrite, martensite, and pearlite is obtained as follows.
First, a plate thickness cross section (L cross section) parallel to the rolling direction of the steel plate is polished. The polished steel plate is corroded using 3% by volume nital. 10 views of the corroded steel sheet thickness ¼ position (position corresponding to ¼ of the sheet thickness in the depth direction from the steel sheet surface) at a magnification of 2000 times using a scanning electron microscope (SEM). Observe and obtain tissue images. About the obtained structure | tissue image, using Image-Pro of Media Cybernetics, the area ratio of each structure | tissue (ferrite, a martensite, pearlite) is calculated for 10 visual fields, and those values are calculated | required and averaged. In the above structure image, ferrite is identified by a gray structure (underlying structure), martensite is a white structure, and pearlite is a layered structure of cementite and ferrite.

<< Volume ratio of retained austenite: 5.0% or less >>
In the high-strength steel sheet of the present invention, in order to sufficiently generate desired martensite and pearlite, it is preferable that the amount of retained austenite is small, and the volume fraction of retained austenite is preferably 5.0% or less, and 3.0% or less. More preferred.
The lower limit of retained austenite is not particularly limited, but the volume ratio of retained austenite is, for example, 0.0% or more.

  The volume ratio of retained austenite is determined by polishing the steel sheet to a thickness of 1/4 surface (surface corresponding to 1/4 of the thickness in the depth direction from the surface of the steel plate), and the diffraction X-ray intensity of this thickness 1/4 surface. Is obtained by measuring MoKα rays are used as incident X-rays, and {111}, {200}, {220}, {311} planes of the retained austenite have peak integrated intensities of ferrite {110}, {200}, {211}. The intensity ratios of all 12 combinations with respect to the integrated intensity of the peak of the surface are obtained, and the average value thereof is taken as the volume ratio of retained austenite.

  The steel structure (microstructure) of the high-strength steel sheet of the present invention may contain carbides such as cementite and bainitic ferrite in addition to ferrite, martensite, pearlite and retained austenite. These structures may be included as long as the total area ratio is in the range of 10.0% or less, and the effects of the present invention are not impaired.

<Value obtained by dividing the amount of Mn in steel by the amount of Mn in ferrite: 1.10 or more>
In the high-strength steel sheet of the present invention, the value obtained by dividing the amount of Mn in steel (unit: mass%) by the amount of Mn in ferrite (unit: mass%) is 1.10 or more.
When holding for a long time such as batch annealing by BAF in the two-phase region of ferrite and austenite, Mn diffuses from the ferrite into the austenite, and the amount of Mn is reduced, that is, a clean ferrite is generated. , Improve ductility and deep drawability. When the value obtained by dividing the amount of Mn in steel by the amount of Mn in ferrite is 1.10 or more, excellent ductility and deep drawability can be secured.
The upper limit of the value obtained by dividing the amount of Mn in steel by the amount of Mn in ferrite is not particularly limited, but is, for example, 3.00 or less from the viewpoint of hole expansibility and chemical conversion treatment.

<Value obtained by dividing the amount of Mn in martensite by the amount of Mn in steel: 1.80 or more>
In the high-strength steel sheet of the present invention, the value obtained by dividing the amount of Mn (unit: mass%) in martensite by the amount of Mn (unit: mass%) in steel is 1.80 or more.
By holding for a long time in the two-phase region of ferrite and austenite in batch annealing by BAF, Mn is concentrated in austenite, and a desired martensite amount can be secured even at an extremely slow cooling rate, and a TS of 590 MPa or more is obtained. . When the value obtained by dividing the amount of Mn in martensite by the amount of Mn in steel is 1.80 or more, a desired TS can be secured.
The upper limit of the value obtained by dividing the amount of Mn in martensite by the amount of Mn in steel is not particularly limited, but is, for example, 4.00 or less from the viewpoint of hole expansibility and chemical conversion property.

  The amount of Mn (unit: mass%) in steel is the amount of Mn (unit: mass%) in the above-described component composition.

The amount of Mn in ferrite (unit: mass%) and the amount of Mn in martensite (unit: mass%) are determined as follows.
First, an EPMA (Electron Probe Micro Analyzer) is used to quantify the distribution state of Mn to each phase of the cross section in the rolling direction at the 1/4 position of the plate thickness. Next, the amount of Mn of 30 ferrite grains and 30 martensite grains is analyzed, and the amount of Mn of each ferrite grain and martensite grain obtained from the analysis result is averaged.

<P (222) / {P (200) + P (220)} ≧ 2.0>
The high-strength steel sheet of the present invention preferably satisfies the following formula in order to ensure better deep drawability.
P (222) / {P (200) + P (220)} ≧ 2.0
In the above formula, P (222), P (200), and P (220) are respectively the (222) plane, the (200) plane, and the (220) plane parallel to the plate plane at the ¼ plate thickness position. Represents the diffraction X-ray integral intensity ratio. As described above, a high intensity ratio of the (222) plane corresponds to the development of the {111} texture.

  The diffraction X-ray integral intensity ratio is a relative intensity when the diffraction X-ray integral intensity of a non-directional standard sample (irregular sample) is used as a reference. The X-ray diffraction measurement may be either an angular dispersion type or an energy dispersion type, and the X-ray source may be a characteristic X-ray or a white X-ray. As measurement surfaces, it is desirable to measure 7 to 10 surfaces (110) to (420) which are the main diffraction surfaces of α-Fe. The steel plate 1/4 plate thickness position refers to a range of 1/8 to 3/8 of the plate thickness of the steel plate as measured from the steel plate surface. The X-ray diffraction measurement may be performed on any surface within this range.

[Production method of low yield ratio type high strength steel sheet]
Next, the manufacturing method of the low yield ratio type high strength steel sheet of the present invention (hereinafter, also simply referred to as “the manufacturing method of the present invention”) will be described.

The production method of the present invention is a method for producing the above-described high-strength steel plate of the present invention. More specifically, a hot-rolled steel sheet is obtained by subjecting a steel slab having the above-described component composition to hot rolling at a finish rolling exit temperature of 800 ° C. or higher and 1000 ° C. or lower. The hot rolled steel sheet is cooled at an average cooling rate of 20 ° C./second to 120 ° C./second in the temperature range from the finish rolling exit temperature to 650 ° C., and the temperature range from 650 ° C. to the following winding temperature is set. Cool at an average cooling rate of 5 ° C./second or more and 40 ° C./second or less. The hot rolled steel sheet after the cooling is wound at a winding temperature of 400 ° C. or higher and 600 ° C. or lower. The hot-rolled steel sheet wound up is pickled. The cold rolled steel sheet is obtained by subjecting the hot rolled steel sheet subjected to the pickling to cold rolling at a rolling reduction of 30% or more and 85% or less. The cold-rolled steel sheet is subjected to batch annealing that satisfies the following temperature histories 1 and 2. By subjecting the cold-rolled steel sheet subjected to the batch annealing to temper rolling at an elongation of 1.1% or less, the above-described high-strength steel sheet of the present invention is obtained.
Temperature history 1: The position of 5 mm along the radial direction from the outer peripheral surface to the inner peripheral surface of the coil of the cold-rolled steel sheet is heated to 620 ° C. or higher and 760 ° C. or lower for 2.0 hours or more and 72.0 or more. The temperature range of 400 ° C. or more and 550 ° C. or less is kept at an average cooling rate of 30 ° C./hour or more and 200 ° C./hour or less.
Temperature history 2: After the temperature was raised from the inner peripheral surface to the outer peripheral surface of the coil of the cold-rolled steel sheet, the position where 1/3 of the coil thickness was entered was 2.0 ° C. to 600 ° C. to 740 ° C. The temperature range of 400 ° C. to 550 ° C. is maintained at an average cooling rate of 5 ° C./hour to 100 ° C./hour.

  Hereinafter, each condition in the manufacturing method of this invention is demonstrated in detail.

<Steel slab>
In the production method of the present invention, a steel slab having the above-described component composition is used.
Steel slabs are usually heated. The heating temperature of the steel slab is preferably 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 the strength. Therefore, the Ti and Nb-based precipitates precipitated during casting are redissolved. Better. When the heating temperature of the steel slab is less than 1100 ° C., it is difficult to sufficiently dissolve the carbide, and there is a possibility that it may lead to troubles during hot rolling due to an increase in rolling load. For this reason, the heating temperature of the steel slab is preferably 1100 ° C. or higher. From the viewpoint of scaling off defects such as bubbles and segregation on the surface of the slab, reducing cracks and irregularities on the steel sheet surface, and achieving a smooth steel sheet surface, the heating temperature of the steel slab is preferably 1100 ° C. or higher.
On the other hand, when the heating temperature of the steel slab is higher than 1300 ° C., the scale loss may increase as the oxidation amount increases. For this reason, the heating temperature of the steel slab is preferably 1300 ° C. or lower.

  The steel slab is preferably manufactured by a continuous casting method in order to prevent macro segregation, but can also be manufactured by an ingot-making method or a thin slab casting method. After manufacturing a steel slab, the conventional method of once cooling to room temperature and heating again after that can also be used. After manufacturing the steel slab, do not cool to room temperature, insert it into a heating furnace as it is, or perform energy saving processes such as direct feed rolling and direct rolling that immediately roll after keeping a little heat Applicable. The steel slab is made into a sheet bar by rough rolling under normal conditions. When the heating temperature is lowered, it is preferable to heat the sheet bar using a bar heater or the like before finish rolling from the viewpoint of preventing troubles during hot rolling.

<Hot rolling finish rolling delivery temperature: 800 ° C or higher and 1000 ° C or lower>
A hot-rolled steel sheet is obtained by subjecting a steel slab (steel slab after heating) to hot rolling including rough rolling and finish rolling.
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, pickling, and the steel sheet after cold rolling. The surface quality tends to deteriorate, and the chemical conversion processability decreases. Furthermore, since the crystal grain size becomes excessively large, a desired martensite amount cannot be obtained, and it becomes difficult to secure TS of 590 MPa or more. There is also a possibility that surface roughness of the pressed product may occur during processing.
On the other hand, when the finish rolling outlet temperature is less than 800 ° C., the rolling load increases, the rolling load increases, or the rolling reduction in a state where austenite is not recrystallized becomes high. As a result, an abnormal texture develops, making it difficult to develop a {111} recrystallized texture of the ferrite of the final material (after annealing), and the deep drawability decreases.
For this reason, the finish rolling exit temperature of hot rolling is 800 ° C. or higher and 1000 ° C. or lower, and preferably 850 ° C. or higher and 950 ° C. or lower.

  During hot rolling, rough rolled sheets may be joined together and finish rolled continuously. The rough rolled plate may be wound up once. In order to reduce the rolling load during hot rolling, part or all of the finish rolling may be lubricated rolling. Lubrication rolling is also effective from the viewpoint of uniform steel plate shape and uniform material. The friction coefficient during lubrication rolling is preferably 0.10 or more and 0.25 or less.

  The hot-rolled steel sheet thus obtained is cooled (primary cooling and secondary cooling) as described below.

<Average cooling rate of primary cooling (cooling in the temperature range from finish rolling delivery temperature to 650 ° C.): 20 ° C./second or more and 120 ° C./second or less>
The cooling (primary cooling) in the temperature range from the finish rolling exit temperature to over 650 ° C. will be described.
When the average cooling rate of primary cooling is less than 20 ° C./second, ferrite transformation and pearlite transformation proceed excessively, and the desired area ratio of bainite cannot be obtained in the structure of the hot-rolled steel sheet after winding. As a result, in a high-strength steel sheet obtained through batch annealing, it becomes difficult to secure a TS of 590 MPa or more at Cold point described later.
On the other hand, when the average cooling rate of primary cooling exceeds 120 ° C./second, a shape failure of the steel sheet occurs.
For this reason, the average cooling rate of primary cooling is 20 ° C./second or more and 120 ° C./second or less, and preferably 25 ° C./second or more and 100 ° C./second or less.

<Average cooling rate of secondary cooling (cooling in the temperature range from 650 ° C. to winding temperature): 5 ° C./second or more and 40 ° C./second or less>
Secondary cooling is performed after the primary cooling. Secondary cooling is cooling in the temperature range from 650 ° C. to the coiling temperature.
When the average cooling rate of the secondary cooling is less than 5 ° C./second, ferrite transformation and pearlite transformation proceed excessively, and a desired area ratio of bainite cannot be obtained in the structure of the hot-rolled steel sheet after winding. As a result, in a high-strength steel sheet obtained through batch annealing, it becomes difficult to secure a TS of 590 MPa or more at Cold point described later.
On the other hand, when the average cooling rate of the secondary cooling exceeds 40 ° C./second, a shape failure of the steel sheet occurs.
For this reason, the average cooling rate of the secondary cooling is 5 ° C./second or more and 40 ° C./second or less, and preferably 5 ° C./second or more and 35 ° C./second or less.

<Taking-up temperature: 400 to 600 ° C.>
The hot rolled steel sheet after cooling (primary cooling and secondary cooling) is wound up.
At this time, if the coiling temperature exceeds 600 ° C., the ferrite transformation or pearlite transformation proceeds excessively, and the desired area ratio of bainite cannot be obtained in the structure of the hot-rolled steel sheet after winding. As a result, in a high-strength steel sheet obtained through batch annealing, it becomes difficult to secure a TS of 590 MPa or more at Cold point described later.
On the other hand, when the coiling temperature is less than 400 ° C., a structure mainly composed of martensite is obtained, the strength is significantly increased, the rolling load in cold rolling is increased, and a steel plate shape defect occurs. .
For this reason, winding temperature is 400 degreeC or more and 600 degrees C or less, and 450 degreeC or more and 600 degrees C or less are preferable.

The wound hot-rolled steel sheet may be referred to as a “material hot-rolled steel sheet”.
In order to secure TS after batch annealing and narrow the TS variation in the coil, the structure of the material hot-rolled steel sheet is preferably mainly composed of bainite, and more specifically, the area ratio of bainite is 70.0. It is preferable that it is more than%. The remaining structure is, for example, polygonal ferrite, pearlite, or carbide.

<Pickling>
The wound hot-rolled steel sheet (raw material hot-rolled steel sheet) is pickled. Since pickling can remove oxides (scale) on the surface of the steel sheet, it is important for ensuring good chemical conversion treatment and plating quality of the high-strength steel sheet as the final product. Pickling may be performed only once or in multiple steps. Pickling may be performed according to a conventional method.

<Cold rolling reduction ratio: 30% to 85%>
Cold rolling is performed on the hot-rolled steel sheet after pickling. Thereby, a cold-rolled steel sheet is obtained.
At this time, if the rolling reduction of cold rolling is less than 30%, the {111} recrystallized texture of ferrite does not sufficiently develop, and excellent deep drawability cannot be obtained. On the other hand, if the rolling reduction of cold rolling exceeds 85%, the load in cold rolling increases, and there is a possibility that a sheet trouble will occur.
For this reason, the rolling reduction of cold rolling is 30% or more and 85% or less, and preferably 40% or more and 75% or less.

Lubricating oil may adhere to the surface of the cold-rolled steel sheet obtained by cold rolling. Oily substances such as lubricating oil may be decomposed during batch annealing by BAF, which will be described later, and remain on the surface of the cold-rolled steel sheet, thereby reducing the surface quality.
Therefore, it is preferable to perform degreasing to remove the lubricating oil on the surface of the cold-rolled steel sheet, for example, in an electrolytic cleaning line, after cold rolling and before performing batch annealing by BAF described later.

<Batch annealing>
Next, batch annealing is performed on the cold-rolled steel sheet. More specifically, the cold rolled steel sheet is wound into a coil shape, and batch annealing is performed on the coil of the cold rolled steel sheet by BAF. At this time, batch annealing is performed so that Hot point (HP) and Cold point (CP) described later in the coil of the cold-rolled steel sheet satisfy a specific temperature history.

The batch annealing by BAF will be described in more detail.
First, before batch annealing, coils of a plurality of stages of cold-rolled steel sheets are arranged in the BAF. The configuration of the BAF and the arrangement of the coils of the cold rolled steel sheet are not particularly limited.
Thereafter, while raising the temperature in the BAF, an atmospheric gas is supplied into the BAF and batch annealing is performed.
The atmosphere gas in the BAF is not particularly limited, and examples thereof include a mixed gas obtained by mixing nitrogen (N ₂ ) gas and hydrogen (H ₂ ) gas, or 100% by volume H ₂ gas.
Among these, H ₂ gas has higher thermal conductivity than N ₂ gas, so the heating rate and cooling rate are increased, productivity is improved, and uniform heating (material stability) and good surface quality are achieved. Therefore, as the atmospheric gas in the BAF, 100% by volume of H ₂ gas is preferable.

FIG. 1 is a perspective view showing a coil 1 with a part cut away. A coil 1 shown in FIG. 1 is formed of a cold-rolled steel sheet wound in a cylindrical shape. The coil 1 has an outer peripheral surface 2 and an inner peripheral surface 3 that are formed by a rolled cold-rolled steel plate.
Hot point (HP) is a position 5 mm along the radial direction r from the outer peripheral surface 2 to the inner peripheral surface 3 of the coil 1.
Cold point (CP) is a position where 1/3 of the coil thickness t is entered along the radial direction r from the inner peripheral surface 3 to the outer peripheral surface 2 of the coil 1.
Both HP and CP are the center positions of the plate width W of the coil 1.

<< Temperature History 1: Temperature History of Hot Point (HP) >>
First, the temperature of the coil of the cold rolled steel sheet is raised. The temperature raising conditions are not particularly limited, but it is preferable to keep HP in a temperature range of 250 ° C. or higher and 450 ° C. or lower for 0.5 hours or more and 8.0 hours or less during temperature rising. Thereby, not only the moisture adhering to the steel plate surface but also the moisture generated by the reduction of the iron oxide on the steel plate surface can be effectively excluded, and the atmospheric gas dew point can be lowered in a very short time. As a result, it is possible to prevent occurrence of temper color, which is uneven color tone, and to improve chemical conversion processability.

(Holding at 620 ° C. to 760 ° C. for 2.0 hours to 72.0 hours)
After the temperature rise, the HP of the coil of the cold rolled steel sheet is held (soaked) at 620 ° C. or higher and 760 ° C. or lower for 2.0 hours or longer and 72.0 hours or shorter.
When the soaking temperature of HP is less than 620 ° C. or the soaking time of HP is less than 2.0 hours, sufficient austenite is not generated, and finally cementite remains undissolved, and the desired martensite and pearlite cannot be secured. Not only is it difficult to secure a TS of 590 MPa or more, but hole expansibility is also reduced.
When the soaking temperature of HP exceeds 760 ° C., the amount of austenite during annealing increases, the amount of Mn concentration in the austenite decreases, and ferrite transformation and bainite transformation progress during cooling after annealing, and the desired martensite. The amount cannot be obtained, and it becomes difficult to secure TS of 590 MPa or more. Ferrite is not sufficiently cleaned, and ductility and deep drawability are also reduced. Furthermore, the concentration of Mn on the surface of the steel sheet during annealing is promoted, and the chemical conversion processability is also lowered.
When the soaking time of HP exceeds 72.0 hours, not only the amount of Mn enriched in austenite is saturated, but also the cost increases. Concentration of Mn on the surface of the steel sheet during annealing is promoted, and chemical conversion processability is also lowered.
The soaking temperature of HP is preferably 640 ° C or higher and 740 ° C or lower. The soaking time of HP is preferably 5.0 hours or more and 60.0 hours or less.

(Average cooling rate in the temperature range from 400 ° C. to 550 ° C .: 30 ° C./hour to 200 ° C./hour)
After the holding, the coil of the cold rolled steel sheet is cooled. At this time, in the HP of the coil of the cold rolled steel sheet, a temperature range of 400 ° C. or more and 550 ° C. or less is cooled at an average cooling rate of 30 ° C./hour or more and 200 ° C./hour or less.
When the average cooling rate in the above temperature range of HP is less than 30 ° C./hour, ferrite transformation and bainite transformation proceed, and the desired amount of martensite cannot be obtained, making it difficult to secure TS of 590 MPa or more.
On the other hand, when the average cooling rate in the above temperature range of HP exceeds 200 ° C./hour, the desired amount of pearlite cannot be obtained, and the hole expandability is lowered.

<< Temperature History 2: Cold Point (CP) Temperature History >>
The temperature of the coil of the cold-rolled steel sheet is raised, and thereafter, the CP of the coil of the cold-rolled steel sheet is held and cooled as follows.

(Holding at 600 ° C to 740 ° C for 2.0 hours to 72.0 hours)
After the temperature rise, the CP of the coil of the cold-rolled steel sheet is held (soaked) at 600 ° C. or higher and 740 ° C. or lower for 2.0 hours or longer and 72.0 hours or shorter.
When the soaking temperature of CP is less than 600 ° C. or the soaking time of CP is less than 2.0 hours, sufficient austenite is not generated, and finally cementite remains undissolved, and the desired martensite and pearlite cannot be secured, Not only is it difficult to secure a TS of 590 MPa or more, but hole expansibility is also reduced.
When the soaking temperature of CP exceeds 740 ° C., the amount of austenite during annealing increases, the amount of Mn concentration in the austenite decreases, and ferrite transformation and bainite transformation proceed during cooling after annealing, and the desired martensite. The amount cannot be obtained, and it becomes difficult to secure TS of 590 MPa or more. Ferrite is not sufficiently cleaned, and ductility and deep drawability are also reduced. Furthermore, the concentration of Mn on the surface of the steel sheet during annealing is promoted, and the chemical conversion processability is also lowered.
When the soaking time of CP exceeds 72.0 hours, not only the amount of Mn enriched in austenite is saturated, but also the cost increases. Concentration of Mn on the surface of the steel sheet during annealing is promoted, and chemical conversion processability is also lowered.
The soaking temperature of CP is preferably 620 ° C. or higher and 720 ° C. or lower. The soaking time of CP is preferably 5.0 hours or more and 60.0 hours or less.

(Average cooling rate in the temperature range from 400 ° C. to 550 ° C .: 5 ° C./hour to 100 ° C./hour)
After the holding, the coil of the cold rolled steel sheet is cooled. At this time, in the CP of the coil of the cold rolled steel sheet, a temperature range of 400 ° C. or more and 550 ° C. or less is cooled at an average cooling rate of 5 ° C./hour or more and 100 ° C./hour or less.
When the average cooling rate of CP in the above temperature range is less than 5 ° C./hour, ferrite transformation and bainite transformation proceed, and a desired martensite amount cannot be obtained, making it difficult to secure TS of 590 MPa or more.
On the other hand, when the average cooling rate of the above temperature range of CP exceeds 100 ° C./hour, a desired amount of pearlite cannot be obtained, and the hole expandability is lowered.

<Elongation rate of temper rolling: 1.1% or less>
The cold-rolled steel sheet subjected to batch annealing is subjected to tempering (skin pass) rolling at an elongation rate of 1.1% or less. When the elongation of temper rolling is more than 1.1%, it is difficult to secure a YR of 0.70 or less. For this reason, the elongation rate of temper rolling is 1.1% or less, and preferably 0.9% or less.
The temper rolling is also effective in correcting coil curl caused by batch annealing by BAF, correcting the shape of the steel sheet, and adjusting the surface roughness. For this reason, the elongation of temper rolling is preferably 0.2% or more.
The temper rolling is performed on the temper rolling line. The temper rolling at the desired elongation rate may be performed at once, or may be performed in several steps.

  In this way, the high strength steel plate of the present invention is manufactured. The manufactured high-strength steel sheet of the present invention may be pickled by a conventional method, and may be subjected to various painting treatments such as resin or oil coating.

  Hereinafter, the present invention will be specifically described with reference to examples. However, the present invention is not limited to these.

<Manufacture of high-strength steel plates (steel plates)>
Steel having the composition shown in the following Table 1 (the balance consisting of Fe and inevitable impurities) was melted in a converter, and a steel slab was obtained by a continuous casting method.
The obtained steel slab was hot-rolled under the conditions shown in Tables 2 to 3 below to obtain hot-rolled steel sheets. The obtained hot-rolled steel sheet was cooled (primary cooling and secondary cooling) and wound under the conditions shown in the same table. The picked hot-rolled steel sheet (raw material hot-rolled steel sheet) was pickled and then cold-rolled under the conditions shown in the same table to obtain a cold-rolled steel sheet. Thereafter, the obtained cold-rolled steel sheet was degreased in an electrolytic cleaning line, and then subjected to batch annealing with BAF under the conditions shown in the same table. Next, temper rolling was performed under the conditions shown in the same table to obtain a high-strength steel plate (steel plate).
About the obtained steel plate, the steel structure (micro structure) was investigated by the method mentioned above. The results are shown in Tables 4 to 7 below.

<Evaluation of high-strength steel plate (steel plate)>
Various characteristics of the obtained high strength steel plate (steel plate) were evaluated. Specifically, the tensile test, the average r value measurement, the deep drawing test, the hole expansion test, the Young's modulus measurement, and the chemical conversion treatment described below were evaluated to evaluate various properties. The results are shown in Tables 4 to 7 below.

<Tensile test>
From the obtained steel plate, a JIS No. 5 test piece was sampled so that the tensile direction was perpendicular to the rolling direction of the steel plate. Using the collected specimens, a tensile test is performed according to JIS Z 2241 (2011), and YP (yield stress), YR (yield ratio), TS (tensile strength), and EL (total elongation) are measured. did. YR is a value obtained by dividing YP by TS.
When YR ≦ 0.70, it was determined to be a low yield ratio type.
It was determined that the strength was high when TS ≧ 590 MPa.
In the case of TS: 590 MPa class, EL ≧ 26%, and in the TS: 780 MPa class, EL ≧ 21%, it was determined that the ductility was good.
TS: 590 MPa class means that TS is 590 MPa or more and less than 780 MPa, and TS: 780 MPa class means that TS is 780 MPa or more and less than 900 MPa.
Further, an intensity difference ΔTS between HP and CP was obtained from the following formula. When ΔTS ≦ 60 MPa, it was determined that there was little strength variation in the coil and the material stability was good.
ΔTS = | TS (HP) −TS (CP) |
In the above formula, TS (HP) represents the tensile strength (unit: MPa) of HP, and TS (CP) represents the tensile strength (unit: MPa) of CP.

<< Measurement of average r value >>
The average r value is an index of deep drawability. The higher the average r value, the better the deep drawability.
The average r value was measured as follows. First, a JIS No. 5 test piece defined in JIS Z 2201 (1998) was collected from the steel plate. More specifically, from the three directions of the rolling direction (L direction) of the steel plate, the 45 ° direction (D direction) with respect to the rolling direction of the steel plate, and the direction perpendicular to the rolling direction of the steel plate (C direction), JIS5 No. test specimens were collected. Using the collected specimen, 10% plastic strain was applied in accordance with JIS Z 2254, and the plastic strain ratio r (r _L , r _D , and r _C ) was determined. The value was calculated. r _L is the plastic strain ratio in the L direction, r _D is the plastic strain ratio in the D direction, and r _C is the plastic strain ratio in the C direction.
Average r value = (r _L + 2r _D + r _C ) / 4
When the average r value ≧ 1.05, it was determined that the deep drawability was good.

<< Deep drawing molding test >>
As the deep drawing test, a cylindrical drawing test was performed, and the deep drawing property was evaluated by the limit drawing ratio (LDR). In the cylindrical deep drawing test, a cylindrical punch having a diameter of 33 mmφ was used. For example, a die having a plate thickness of 1.2 mm and a die having a die diameter of 36.6 mm were used (other thicknesses will be described later). The test was performed with a wrinkle holding force of 1.5 ton (14.71 kN). Since the sliding state of the surface changes depending on the plating state or the like, the test was performed under high lubrication conditions by placing a polyethylene sheet between the sample and the die so that the sliding state of the surface did not affect the test. The blank diameter was changed at a pitch of 1 mm, and the ratio (D / d) of the blank diameter D and punch diameter d that was not ruptured and squeezed out was defined as LDR.
When LDR ≧ 2.00, it was determined that the deep drawability was good.

The die diameter of the mold used for the deep drawing test (cylindrical drawing test) was as follows for each plate thickness of the steel plate.
-Thickness 0.8mm Material ... Die diameter of mold: 35.4mm
・ Thickness of 1.0mm material ... Die diameter of mold: 36.0mm
・ Thickness of 1.2mm material ・ ・ ・ Die diameter of mold: 36.6mm
・ Thickness of 1.4mm material ・ ・ ・ Die diameter of mold: 37.2mm
-Material with a plate thickness of 1.6 mm ... Die diameter of the mold: 37.8 mm
・ Thickness 1.8mm material ・ ・ ・ Die diameter of mold: 38.4mm
・ Thickness 2.0mm material ・ ・ ・ Die diameter of mold: 39.0mm
・ Thickness 2.3mm Material ... Die diameter of mold: 39.9mm

《Hole expansion test》
The hole expansion test was conducted in accordance with JIS Z 2256 (2010). More specifically, first, the obtained steel plate was cut into 100 mm × 100 mm. A 10 mm diameter hole was punched into the cut steel sheet with a clearance of 12% ± 1%. Thereafter, using a die having an inner diameter of 75 mm, a punch having a 60 ° conical shape was pushed into the hole while being suppressed with a wrinkle holding force of 9 ton (88.26 kN), and the hole diameter at the crack initiation limit was measured. Then, the critical hole expansion ratio λ [%] was obtained from the following formula.
Limit hole expansion ratio λ [%] = {(D _f −D ₀ ) / D ₀ } × 100
However, _{D f} is the hole diameter at crack initiation [mm], _{D 0} is the initial hole diameter [mm].
In the case of λ ≧ 50% in the TS590 MPa class and in the case of λ ≧ 30% in the TS780 MPa class, it was determined that the hole expandability was good.

<Measurement of Young's modulus>
The Young's modulus is an index of rigidity and improves as the {111} recrystallized texture of ferrite develops.
About the obtained steel plate, the test piece of 10 mm x 50 mm was extract | collected from the orthogonal | vertical direction (C direction) with respect to the rolling direction. For the collected specimens, Young's modulus was measured using a transverse vibration type resonance frequency measuring device in accordance with American Society to Testing Materials standards (C1259).
When the Young's modulus was 220 GPa or more, it was determined that the rigidity was good.

<Method for evaluating chemical conversion treatment>
The obtained steel sheet was subjected to a chemical conversion treatment by the following method using a chemical conversion treatment liquid (Palbond L3080 (registered trademark)) manufactured by Nihon Parkerizing Co., Ltd., and a chemical conversion treatment property was evaluated.
First, the obtained steel sheet was degreased using a degreasing liquid Fine Cleaner (registered trademark) manufactured by Nihon Parkerizing Co., Ltd., then washed with water, and then used with a surface conditioning liquid preparen Z (registered trademark) manufactured by Nihon Parkerizing Co., Ltd. For 30 seconds. The surface-adjusted steel sheet was immersed in a chemical conversion treatment solution (Palbond L3080) at 43 ° C. for 120 seconds, then washed with water and dried with warm air. Thus, the steel sheet was subjected to chemical conversion treatment.
About the surface of the steel plate after chemical conversion treatment, 5 visual fields were observed at random by 500 times using SEM (scanning electron microscope). The area ratio [%] of the region (skee) where the chemical conversion film was not generated was obtained by image processing, and the following evaluation was performed based on the obtained area ratio.
Score 5: 5% or less Score 4: More than 5% but 10% or less Score 3: More than 10% and 25% or less Score 2: More than 25% and 40% or less Score 1: More than 40% Chemical conversion processability if score 4 or score 5 Was determined to be good.

In Table 1 to Table 7, the underlined portion indicates the outside of the scope of the present invention or the outside of the preferred range.
In Tables 4 to 7, F represents ferrite, M represents martensite, P represents pearlite, RA represents retained austenite, θ represents carbide (TiC, NbC, cementite, etc.), and B represents bainitic ferrite.

From the results shown in Tables 1 to 7, the steel sheet of the example of the present invention has a low yield ratio (YR) of 0.70 or less and a tensile strength (TS) of 590 MPa or more. The properties (ductility, deep drawability, hole expansibility) and chemical conversion treatment were good.
On the other hand, the steel plate of the comparative example has insufficient one or more of the above characteristics.

  When the steel plates of the present invention are compared, the above formula is satisfied rather than the case where the formula: P (222) / {P (200) + P (220)} ≧ 2.0 is not satisfied (No. 3 and No. 7). In the case, the deep drawability (average r value and LDR) was more excellent.

1: Coil 2: Outer peripheral surface 3: Inner peripheral surface CP: Cold Point
HP: Hot Point
r: Radial direction of coil t: Coil thickness W: Coil plate width

Claims (5)

By mass%, C: more than 0.050% and less than 0.200%, Si: 0.01% to less than 0.50%, Mn: more than 1.80% and less than 2.60%, P: 0.100% or less , S: 0.0200% or less, Al: 2.000% or less, and N: 0.0100% or less, and Ti: 0.005% or more and 0.100% or less, and Nb: 0 A composition comprising at least one selected from the group consisting of 0.005% or more and 0.100% or less, with the balance being Fe and inevitable impurities;
A steel structure having an area ratio of ferrite of 40.0% to 90.0%, martensite of 5.0% to 30.0%, and pearlite of 2.0% to 30.0%. Have
The value obtained by dividing the amount of Mn in steel by the amount of Mn in the ferrite is 1.10 or more, and the value obtained by dividing the amount of Mn in the martensite by the amount of Mn in steel is 1.80 or more. Low yield ratio type high-strength steel sheet, the unit of which is mass%. The low yield ratio type high-strength steel sheet according to claim 1, wherein the component composition further contains at least one selected from the following groups A to E in mass%.
[Group A] Ni: 0.01% or more and 1.00% or less, Cu: 0.005% or more and 1.000% or less, Cr: 0.01% or more and 1.00% or less, and Mo: 0.005 %: At least one selected from the group consisting of 0.5% to 0.500% [Group B] V: 0.005% to 0.100% and W: 0.005% to 0.100% [Group C] Ca: 0.0005% or more and 0.0050% or less, Mg: 0.0005% or more and 0.0050% or less, and REM: 0.0005% or more and 0.0050% or less At least one selected from the group consisting of [Group D] Sb: 0.002% to 0.200%, Sn: 0.002% to 0.200%, and Ta: 0.001% to 0.000. Selected from the group consisting of 010% or less At least one [E group] B: 0.0001% or more 0.0050% or less Furthermore, the low yield ratio type | mold high strength steel plate of Claim 1 or 2 which satisfy | fills a following formula.
P (222) / {P (200) + P (220)} ≧ 2.0
In the above formula, P (222), P (200) and P (220) are respectively the (222) plane, the (200) plane and the (220) plane parallel to the plate plane at the steel plate 1/4 plate thickness position. Represents the diffraction X-ray integral intensity ratio.   The low yield ratio type high-strength steel sheet according to any one of claims 1 to 3, wherein a volume ratio of retained austenite in the steel structure is 5.0% or less. A hot-rolled steel sheet is obtained by subjecting the steel slab having the composition according to claim 1 or 2 to hot rolling at a finish rolling exit temperature of 800 ° C. or higher and 1000 ° C. or lower,
The hot rolled steel sheet is cooled at an average cooling rate of 20 ° C./second to 120 ° C./second in a temperature range from the finish rolling outlet temperature to 650 ° C., and a temperature range from 650 ° C. to the following winding temperature is set. Cool at an average cooling rate of 5 ° C / second or more and 40 ° C / second or less,
Winding the hot-rolled steel sheet after the cooling at a winding temperature of 400 ° C. or higher and 600 ° C. or lower,
The picked hot-rolled steel sheet is pickled,
By cold-rolling the hot-rolled steel sheet subjected to the pickling at a reduction rate of 30% to 85%, a cold-rolled steel sheet is obtained,
The cold-rolled steel sheet is subjected to batch annealing satisfying the following temperature history 1 and 2,
The low yield ratio type high-strength steel sheet according to any one of claims 1 to 4, wherein the cold-rolled steel sheet subjected to the batch annealing is subjected to temper rolling at an elongation of 1.1% or less. A method for producing a low yield ratio type high strength steel sheet.
Temperature history 1: A position of 5 mm along the radial direction from the outer peripheral surface to the inner peripheral surface of the coil of the cold-rolled steel sheet is heated to 620 ° C. or higher and 760 ° C. or lower for 2.0 hours or more and 72.0 or more. The temperature range of 400 ° C. or more and 550 ° C. or less is kept at an average cooling rate of 30 ° C./hour or more and 200 ° C./hour or less.
Temperature history 2: After the temperature was raised from 1 to 3rd of the coil thickness along the radial direction from the inner peripheral surface to the outer peripheral surface of the coil of the cold-rolled steel sheet, 2.0 ° The temperature range of 400 ° C. to 550 ° C. is maintained at an average cooling rate of 5 ° C./hour to 100 ° C./hour.