JP6597811B2 - High-strength cold-rolled steel sheet and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet and a method for producing the same. More specifically, the present invention relates to a high-strength cold-rolled steel sheet having a high tensile strength (TS) of 1320 MPa or more and a method for producing the same suitable for parts of transportation machinery including automobiles.

Conventionally, high-strength cold-rolled steel sheets have been applied to body parts and the like (for example, Patent Documents 1 and 2).
In recent years, from the viewpoint of preservation of the global environment, improvement in fuel efficiency of automobiles has been demanded, and application of high-strength cold-rolled steel sheets having a tensile strength of 1320 MPa or more has been promoted.
Furthermore, recently, the demand for improving the collision safety of automobiles has increased, and from the viewpoint of ensuring the safety of passengers in the event of a collision, the tensile strength for structural members such as a skeleton part of a vehicle body is extremely high, being 1460 MPa or more. Application of high-strength cold-rolled steel sheets having strength has also been studied.

International Publication No. 2016/132680 International Publication No. 2016/021193

The ductility decreases as the strength of the steel plate increases. Since a steel sheet with low ductility is cracked during press molding, it is necessary to combine high ductility with high strength in order to process a high-strength steel sheet as an automobile part.
Further, the mechanical properties of the steel sheet are generally anisotropic, and the characteristic values differ depending on the test direction. When the ductility of the steel sheet has large anisotropy, cracks are likely to occur when subjected to a process that extends in a direction with low ductility.
Therefore, a steel sheet having a high tensile strength of 1320 MPa or more, maintaining excellent ductility, and further suppressing ductility anisotropy as much as possible is required.
However, conventional cold-rolled steel sheets may have any of the above characteristics insufficient.

  Accordingly, an object of the present invention is to provide a high-strength cold-rolled steel sheet having a tensile strength of 1320 MPa or more, excellent ductility, and small ductility anisotropy, and a method for producing the same. .

The present inventors have intensively studied to achieve the above object. As a result, in order to achieve both high strength and high ductility, it is usually effective to utilize the TRIP (Transformation Induced Plasticity) phenomenon associated with the processing-induced transformation of retained austenite. In the case where a large amount of acicular residual austenite having a ratio of 0.6 or less is present in an area ratio, it has been found that particularly excellent ductility is exhibited.
Based on the above findings, the present inventors have conducted intensive studies on methods that can further reduce the ductility anisotropy. As a result, in the steel sheet having a microstructure in which acicular residual austenite having an aspect ratio of 0.6 or less is present in a large amount in area ratio, and the location of such residual austenite is mainly a Bain group boundary. Found that the anisotropy of ductility was significantly reduced.
The inventors have further studied. As a result, the heat treatment (annealing process) of the steel sheet was performed twice, and in particular, by optimizing the thermal history in the first annealing process, it was found that the microstructure of the steel sheet can be stably made the above microstructure. .
Based on the above findings, the present inventors have further studied and completed the present invention.

That is, the present invention provides the following [1] to [5].
[1] By mass%, C: more than 0.18% to 0.45% or less, Si: 0.50% to 2.50%, Mn: more than 2.50% to 4.00%, P: 0.00. 050% or less, S: 0.0100% or less, Al: 0.010% or more and 0.100% or less, and N: 0.0100% or less, having a composition consisting of the balance Fe and inevitable impurities, C and Mn of the above composition satisfy the following formula (1) in mass%, and in the microstructure, the total area ratio of ferrite and bainitic ferrite is 10% or more and 50% or less, and the area of residual austenite The ratio of the martensite area ratio of more than 15% to 50% or less and the martensite area ratio of more than 15% to 60% or less of the retained austenite having an aspect ratio of 0.6 or less is 70% in terms of area ratio. % And more Of the residual austenite ratio is 0.6 or less, the percentage of those present in Bain group boundary is an area ratio of 50% or more, high strength cold rolled steel sheet.
7.5 × C + Mn ≧ 5.0 (1)
However, in Formula (1), C and Mn show content of each element.
[2] The above composition is further mass%, Ti: 0.005% to 0.035%, Nb: 0.005% to 0.035%, V: 0.005% to 0.035% Hereinafter, Mo: 0.005% to 0.035%, B: 0.0003% to 0.0100%, Cr: 0.05% to 1.00%, Ni: 0.05% to 1. 00% or less, Cu: 0.05% or more and 1.00% or less, Sb: 0.002% or more and 0.050% or less, Sn: 0.002% or more and 0.050% or less, Ca: 0.0005% or more [50] containing at least one element selected from the group consisting of 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. ] The high-strength cold-rolled steel sheet according to any one of the above
[3] The high-strength cold-rolled steel sheet according to [1] or [2], which has a plating layer on the surface.
[4] A method for producing the high-strength cold-rolled steel sheet according to any one of [1] to [3], wherein the steel material having the composition according to [1] or [2] By rolling, a hot rolling step for obtaining a hot-rolled plate, a pickling step for pickling the hot-rolled plate, and a hot rolling plate subjected to the pickling have a reduction rate of 30% or more. A cold rolling step of obtaining a cold rolled sheet by performing cold rolling, and the cold rolled sheet are heated at an annealing temperature T ₁ of Ac ₃ points or more and 950 ° C. or less, and from the annealing temperature T ₁ to 10 ° C. in / s exceeds the average cooling rate, cooling to cooling stop temperature T ₂ less than 250 ° C. or higher 350 ° C., by holding 10s or above the cooling stop temperature T _2, the first to obtain the first Danhiyanobe annealed sheets a stage annealing step, the first Danhiyanobe annealed sheets, heated at an annealing temperature T ₃ of 680 ° C. or higher 820 ° C. or less, the sintered The temperature T _3, by cooling to 300 ° C. or higher 500 ° C. or less of the cooling stop temperature T _4, the method of producing a high strength cold rolled steel sheet comprising: a second-stage annealing process of obtaining a second Danhiyanobe annealed sheet, a.
[5] The high strength according to [4], further comprising a plating step of subjecting the second-stage cold-rolled annealed plate to a hot dip galvanizing treatment, a hot dip galvanizing treatment and an alloying treatment, or an electrogalvanizing treatment. A method for producing a cold-rolled steel sheet.

According to the present invention, it is possible to provide a high-strength cold-rolled steel sheet having a tensile strength of 1320 MPa or more, excellent ductility, and small ductility anisotropy, and a method for producing the same.
The high-strength cold-rolled steel sheet of the present invention is suitable for structural steel materials such as parts for transportation machinery including automobiles and steel materials for construction. ADVANTAGE OF THE INVENTION According to this invention, a transport equipment member can be made thinner and stronger than before, and further application development of a high-strength cold-rolled steel sheet can be achieved, which has a remarkable industrial effect.

It is a schematic diagram which shows a part (area | region considered to be produced | generated from one prior austenite grain) of the microstructure of a steel plate. The proportion of residual austenite having an aspect ratio of 0.6 or less and the proportion of residual austenite having an aspect ratio of 0.6 or less at the Bain group boundary affect the anisotropy of elongation. It is a graph which shows an influence.

[High-strength cold-rolled steel sheet]
The high-strength cold-rolled steel sheet of the present invention is, in mass%, C: more than 0.18% and 0.45% or less, Si: 0.50% or more and 2.50% or less, Mn: more than 2.50% and 4.00. %: P: 0.050% or less, S: 0.0100% or less, Al: 0.010% or more and 0.100% or less, and N: 0.0100% or less, and the balance Fe and inevitable impurities C and Mn of the above composition satisfy the following formula (1) in mass%, and the total area ratio of ferrite and bainitic ferrite in the microstructure is 10% or more and 50% or less The area ratio of retained austenite is more than 15% and 50% or less, and the area ratio of martensite is more than 15% and 60% or less, and among the retained austenite, the aspect ratio is 0.6 or less Ratio is 70% in area ratio A top, of the residual austenite aspect ratio of 0.6 or less, the percentage of those present in Bain group boundary is 50% or more in area ratio, a high strength cold rolled steel sheet.
7.5 × C + Mn ≧ 5.0 (1)
However, in Formula (1), C and Mn show content of each element.
In addition, the plate | board thickness of the high intensity | strength cold-rolled steel plate of this invention is 5 mm or less, for example.

<composition>
Below, the composition (component composition) which the high-strength cold-rolled steel sheet of this invention has is demonstrated first. The unit of element content in the component composition is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.

<< C: more than 0.18% and less than 0.45% >>
C stabilizes austenite, ensures retained austenite with a desired area ratio, contributes effectively to improving ductility, increases the hardness of martensite, and contributes to an increase in strength. In order to sufficiently obtain such an effect, C needs to contain more than 0.18%.
On the other hand, a large content exceeding 0.45% leads to deterioration of toughness, weldability and delayed fracture resistance, and makes the amount of martensite produced excessively and decreases ductility.
Therefore, the C content is more than 0.18% and 0.45% or less, preferably 0.19% or more and 0.43% or less, and more preferably 0.20% or more and 0.42% or less.

<< Si: 0.50% or more and 2.50% or less >>
Si suppresses the formation of carbide (cementite) and promotes the concentration of C to austenite, thereby stabilizing austenite and contributing to the improvement of the ductility of the steel sheet. Si dissolved in ferrite improves work hardening ability and contributes to improvement of ductility of ferrite itself. In order to sufficiently obtain such an effect, Si needs to be contained in an amount of 0.50% or more.
On the other hand, if Si exceeds 2.50%, not only the effect of suppressing the formation of carbide (cementite) and stabilizing the retained austenite is saturated, but the amount of Si dissolved in ferrite becomes excessive. Therefore, ductility is reduced.
For this reason, the Si content is 0.50% or more and 2.50% or less, preferably 0.80% or more and 2.40% or less, and more preferably 1.00% or more and 2.30% or less.

<< Mn: more than 2.50% and less than 4.00% >>
Mn is an austenite stabilizing element, and contributes to the improvement of ductility by stabilizing austenite, and also promotes the formation of martensite by increasing the hardenability and contributes to increasing the strength of the steel sheet. In order to sufficiently obtain such an effect, Mn needs to be contained in an amount exceeding 2.50%.
On the other hand, when Mn exceeds 4.00%, martensite is excessively generated and ductility is deteriorated.
For this reason, content of Mn is more than 2.50% and 4.00% or less, and preferably 2.70% or more and 3.50% or less.

<< P: 0.050% or less >>
P is a harmful element that segregates at grain boundaries to reduce elongation, induces cracking during processing, and further degrades impact resistance. Therefore, the P content is 0.050% or less. Preferably it is 0.010% or less.
However, excessive P removal causes an increase in refining time and cost, and therefore the P content is preferably 0.002% or more.

<< S: 0.0100% or less >>
S exists as MnS in the steel, promotes the generation of voids during the punching process, and further decreases the stretch flangeability because it becomes a starting point for the generation of voids during the processing. Therefore, the amount of S is preferably reduced as much as possible, and is 0.0100% or less. Preferably it is 0.0050% or less.
However, excessive desulfurization causes an increase in refining time and cost, and therefore the S content is preferably 0.0002% or more.

<< Al: 0.010% or more and 0.100% or less >>
Al is an element that acts as a deoxidizer. In order to obtain such an effect, 0.010% or more of Al is contained.
However, when the Al content is excessive, it remains as an Al oxide in the steel sheet, and the Al oxide is easily aggregated and coarsened, causing a deterioration in local ductility. Therefore, the Al content is set to 0.100% or less.

<< N: 0.0100% or less >>
N is present as AlN in the steel and promotes the generation of coarse voids during the punching process, and further reduces the local ductility because it becomes the starting point for the generation of coarse voids during the machining. For this reason, it is preferable to reduce N amount as much as possible, and N content shall be 0.0100% or less. Preferably it is 0.0060% or less.
However, excessive de-N causes an increase in refining time and an increase in cost, so the N content is preferably 0.0005% or more.

<< 7.5 * C + Mn: 5.0 or more >>
C and Mn are both elements that contribute to increasing the strength of the steel sheet, but 7.5 × C + Mn is less than 5.0 even when the content of each element is individually within the above range. In some cases, the desired steel plate strength cannot be achieved. This is because C and Mn do not independently contribute to an increase in steel sheet strength, but influence each other and form martensite and retained austenite, resulting in an increase in steel sheet strength.
In particular, martensite, which has a strong influence on steel plate strength, changes its contribution to steel plate strength depending on its area ratio, C concentration, and Mn concentration, so it is difficult to control steel plate strength only by the martensite area ratio. However, if 7.5 × C + Mn is 5.0 or more, a desired steel plate strength is achieved.
For this reason, C and Mn must satisfy the following formula (1) in mass%.
7.5 × C + Mn ≧ 5.0 (1)
(However, in formula (1), C and Mn indicate the content of each element.)
7.5 × C + Mn is preferably 5.1 or more.

<Other components (elements)>
In the high-strength cold-rolled steel sheet of the present invention, the above composition is further, if necessary, in mass%, Ti: 0.005% or more and 0.035% or less, Nb: 0.005% or more and 0.035% or less. V: 0.005% to 0.035%, Mo: 0.005% to 0.035%, B: 0.0003% to 0.0100%, Cr: 0.05% to 1.00 %: Ni: 0.05% to 1.00%, Cu: 0.05% to 1.00%, Sb: 0.002% to 0.050%, Sn: 0.002% to 0 .050% or less, 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 element.

(Ti: 0.005% to 0.035%)
Ti forms carbonitrides and increases the strength of the steel by precipitation strengthening action. When adding Ti, in order to exhibit the said effect | action effectively, it is preferable to make Ti content 0.005% or more. On the other hand, if Ti is excessive, precipitates may be generated excessively and ductility may be reduced.
For this reason, the content of Ti is preferably 0.005% or more and 0.035% or less, and more preferably 0.005% or more and 0.020% or less.

(Nb: 0.005% or more and 0.035% or less)
Nb forms carbonitride and increases the strength of the steel by precipitation strengthening action. When Nb is added, the Nb content is preferably 0.005% or more in order to effectively exhibit the above-described action. On the other hand, if Nb is excessive, precipitates may be generated excessively and ductility may be reduced.
For this reason, the Nb content is preferably 0.005% or more and 0.035% or less, and more preferably 0.005% or more and 0.030% or less.

(V: 0.005% to 0.035%)
V forms carbonitride and increases the strength of the steel by precipitation strengthening action. When adding V, in order to exhibit the said effect | action effectively, it is preferable to make V content 0.005% or more. On the other hand, if V is excessive, precipitates may be generated excessively and ductility may be reduced.
For this reason, the content of V is preferably 0.005% or more and 0.035% or less, and more preferably 0.005% or more and 0.030% or less.

(Mo: 0.005% to 0.035%)
Mo forms carbonitrides and increases the strength of the steel by precipitation strengthening action. When adding Mo, in order to exhibit the said effect | action effectively, it is preferable to make Mo content 0.005% or more. On the other hand, if Mo is excessive, precipitates may be generated excessively and ductility may be reduced.
For this reason, the content of Mo is preferably 0.005% or more and 0.035% or less, and more preferably 0.005% or more and 0.030% or less.

(B: 0.0003% or more and 0.0100% or less)
B has an effect of enhancing hardenability and promoting the formation of martensite, and thus is useful as a steel strengthening element. In order to effectively exhibit the above action, the B content is preferably 0.0003% or more. On the other hand, if B is excessive, martensite may be generated excessively and ductility may be reduced.
For this reason, the content of B is preferably 0.0003% or more and 0.0100% or less.

(Cr: 0.05% to 1.00%)
Cr has a function of enhancing hardenability and promoting martensite formation, and thus is useful as a steel strengthening element. In order to effectively exhibit the above action, the Cr content is preferably 0.05% or more. On the other hand, if Cr is excessive, martensite may be generated excessively and ductility may be reduced.
For this reason, the Cr content is preferably 0.05% or more and 1.00% or less.

(Ni: 0.05% or more and 1.00% or less)
Ni is useful as a steel strengthening element because it has the effect of enhancing hardenability and promoting the formation of martensite. In order to effectively exhibit the above action, the Ni content is preferably 0.05% or more. On the other hand, if Ni is excessive, martensite is generated excessively and ductility may be reduced.
For this reason, the Ni content is preferably 0.05% or more and 1.00% or less.

(Cu: 0.05% or more and 1.00% or less)
Cu has a function of enhancing hardenability and promoting the formation of martensite, and thus is useful as a steel strengthening element. In order to effectively exhibit the above action, the Cu content is preferably 0.05% or more. On the other hand, if Cu is excessive, martensite may be generated excessively and ductility may be reduced.
For this reason, the content of Cu is preferably 0.05% or more and 1.00% or less.

(Sb: 0.002% to 0.050%)
Sb has the effect | action which suppresses the decarburization of the steel plate surface layer (area | region of about several tens of micrometers) produced by nitriding and oxidation of the steel plate surface. Thereby, it can prevent that the production amount of austenite reduces on the steel plate surface, and it is effective for ensuring desired ductility. In order to effectively exhibit the above action, the Sb content is preferably 0.002% or more. On the other hand, if Sb is excessive, the toughness may be reduced.
For this reason, the content of Sb is preferably 0.002% or more and 0.050% or less.

(Sn: 0.002% to 0.050%)
Sn has an effect of suppressing decarburization of the steel sheet surface layer (region of about several tens of μm) caused by nitriding and oxidation of the steel sheet surface. Thereby, it can prevent that the production amount of austenite reduces on the steel plate surface, and it is effective for ensuring desired ductility. In order to effectively exhibit the above action, the Sn content is preferably 0.002% or more. On the other hand, if Sn is excessive, the toughness may be reduced.
For this reason, the content of Sn is preferably 0.002% or more and 0.050% or less.

(Ca: 0.0005% or more and 0.0050% or less)
Ca has the effect | action which controls the form of a sulfide type inclusion, and is effective in suppression of the fall of local ductility. When adding Ca, in order to acquire the said effect, it is preferable to make Ca content 0.0005% or more. On the other hand, if the Ca content is excessive, the effect may be saturated.
For this reason, the content of Ca is preferably 0.0005% or more and 0.0050% or less.

(Mg: 0.0005% or more and 0.0050% or less)
Mg has the effect | action which controls the form of a sulfide type inclusion, and is effective in suppression of a local ductility fall. When adding Mg, in order to acquire the said effect, it is preferable to make Mg content 0.0005% or more. On the other hand, if the Mg content is excessive, the effect may be saturated.
For this reason, the content of Mg is preferably 0.0005% or more and 0.0050% or less.

(REM: 0.0005% or more and 0.0050% or less)
REM (rare earth element) has an action of controlling the form of sulfide inclusions, and is effective in suppressing a decrease in local ductility. When adding REM, in order to acquire the said effect, it is preferable to make REM content 0.0005% or more. On the other hand, if the REM content is excessive, the effect may be saturated.
For this reason, the content of REM is preferably 0.0005% or more and 0.0050% or less.

<< Remainder Fe and inevitable impurities >>
In the above composition, the balance other than the above components consists of Fe (remainder Fe) and inevitable impurities.

<Microstructure of steel sheet>
Next, the microstructure in the high-strength cold-rolled steel sheet of the present invention will be described.

<< Total area ratio of ferrite + bainitic ferrite: 10% to 50% >>
Ferrite and bainitic ferrite are soft structures and contribute to improving the ductility of the steel sheet. Since carbon does not dissolve so much in these structures, discharging C into the austenite increases the stability of the austenite and contributes to the improvement of ductility.
In order to impart the required ductility to the steel sheet, the total area ratio of ferrite and bainitic ferrite is required to be 10% or more.
On the other hand, if the sum of the area ratios of ferrite and bainitic ferrite exceeds 50%, it becomes difficult to ensure a tensile strength of 1320 MPa or more.
For this reason, the sum total of the area ratios of ferrite and bainitic ferrite is 10% or more and 50% or less.

<< Area ratio of retained austenite: more than 15% and 50% or less >>
The retained austenite is a structure that is rich in ductility per se, but is a structure that contributes to further improving ductility by strain-induced transformation. In order to obtain such an effect, the retained austenite needs to be more than 15% by area ratio.
On the other hand, if the retained austenite increases in area ratio exceeding 50%, the stability of the retained austenite is lowered, so that strain-induced transformation occurs at an early stage, and ductility is lowered.
For this reason, the area ratio of a retained austenite is more than 15% and 50% or less.
In this specification, the volume ratio of retained austenite is calculated by the method described later, and this is treated as the area ratio.

《Martensite area ratio: more than 15% and less than 60%》
The “martensite” here includes fresh martensite and tempered martensite.
Martensite is a very hard structure and contributes to increasing the strength of the steel sheet. When the martensite content is 15% or less in terms of area ratio, desired steel plate strength cannot be obtained.
On the other hand, when it contains exceeding 60% by area ratio, it becomes impossible to ensure desired ductility.
For this reason, the sum total of the area ratio of martensite is more than 15% and 60% or less. Preferably they are 20% or more and 55% or less.
The microstructure of the high-strength cold-rolled steel sheet according to the present invention is not limited to the case where the total area ratio of each of the ferrite and bainitic ferrite, retained austenite, and martensite is 100%. In some cases, the area ratio is 100%.

<< Ratio of retained austenite having an aspect ratio of 0.6 or less: 70% or more in area ratio >>
Residual austenite improves the ductility of the steel sheet, but its contribution to the improvement of ductility differs depending on its shape. Residual austenite having an aspect ratio of 0.6 or less is more stable to processing than the retained austenite having an aspect ratio exceeding 0.6, and has a greater effect of improving ductility due to strain-induced transformation.
Residual austenite having a low processing stability and an aspect ratio of more than 0.6 becomes hard martensite at an early stage of the tensile test, so that coarse voids are likely to be formed around the austenite. In the latter stage of the tensile test, the coarse voids are easily connected, so that the ductility is lowered.
On the other hand, the retained austenite having an aspect ratio of 0.6 or less is deformed along the flow of the microstructure, and it is difficult to form voids around it.
In addition to these ductility-improving effects, in order to fully enjoy the effect of reducing the anisotropy of elongation due to residual austenite having an aspect ratio of 0.6 or less present at the boundary of the Bain group described later, The ratio of the retained austenite having an aspect ratio of 0.6 or less may be an area ratio of 70% or more. Preferably it is 75% or more.
The upper limit of this ratio is not particularly limited, and may be 100%.

<< Ratio of residual austenite having an aspect ratio of 0.6 or less at the boundary of the Bain group: area ratio of 50% or more >>
First, the retained austenite present at the Bain group boundary will be described.

In martensite and bainite, 24 variants having a Kurdjumov-Sachs (KS) relationship can be generated from one old austenite grain. Variants arising from one old austenite grain are divided into three Bain groups (for example, “Shiroro Miyamoto, 3 others,“ Crystallographic restraint in martensite / bainite transformation of steel ”, Journal of the Japan Institute of Metals, Public Interest Japan Metallurgy Society, July 2015, Vol. 79, No. 7, pp. 339-347)).
Since the high-strength cold-rolled steel sheet of the present invention is obtained through a plurality of annealing steps as will be described later, the microstructure of the steel sheet is different from martensite and bainite transformed from the austenite single phase, but is distinguished from the bcc phase. The same grouping as described above can be performed for the portion to be processed.
FIG. 1 is a schematic diagram showing a part of the microstructure of a steel plate (region considered to be generated from one prior austenite grain). The microstructure of the steel plate shown in FIG. 1 is composed of three Bain groups (B1 to B3). The same Bain group is given the same hatching.
Residual austenite is also present in the microstructure of the steel sheet shown in FIG. The retained austenite indicated by the symbol “RA ₂ ” exists inside one Bain group B2. On the other hand, the retained austenite indicated by the symbol “RA ₁ ” is present at the boundary between the Bain group B1 and another Bain group B3.
The retained austenite indicated by the symbol “RA ₁ ” corresponds to the retained austenite existing at the Bain group boundary.

When residual austenite having an aspect ratio of 0.6 or less is present at the Bain group boundary, the anisotropy of elongation is significantly reduced.
Although this reason is not necessarily clear, the present inventors consider as follows. That is, the strain-induced martensitic transformation behavior of retained austenite usually varies depending on the direction of external stress, but the Bain group boundary has a large misorientation, and stress in various directions as well as in the external force direction is locally applied. concentrate. For this reason, retained austenite existing at the Bain group boundary contributes to the improvement of ductility by causing the same strain-induced martensitic transformation behavior even when the direction of external force is changed. To reduce.
It is the strain-induced martensitic transformation in the latter stage of the tensile test that greatly affects the elongation. For this reason, in order to obtain such an effect, it is necessary that residual austenite having an aspect ratio of 0.6 or less that can exist stably until the latter stage of the tensile test is present at the Bain group boundary.
In order to sufficiently reduce the anisotropy of elongation (ductivity anisotropy), the proportion of residual austenite having an aspect ratio of 0.6 or less at the Bain group boundary is 50% in terms of area ratio. It may be above, preferably 65% or more.
Although the upper limit of this ratio is not specifically limited, For example, the area ratio is 95% or less.

<Plating layer>
The high-strength cold-rolled steel sheet of the present invention may further have a plating layer on the surface from the viewpoint of improving corrosion resistance and the like. As the plating layer, a hot dip galvanized layer, an alloyed hot dip galvanized layer, or an electrogalvanized layer is preferable.
The hot-dip galvanized layer, the alloyed hot-dip galvanized layer, and the electrogalvanized layer are not particularly limited, and are conventionally known hot-dip galvanized layer, conventionally known alloyed hot-dip galvanized layer, and conventionally known, respectively. The electrogalvanized layer is preferably used.
The electrogalvanized layer may be a zinc alloy plated layer obtained by adding an appropriate amount of elements such as Fe, Cr, Ni, Mn, Co, Sn, Pb, or Mo to Zn according to the purpose. .

[Method for producing high-strength cold-rolled steel sheet]
Next, a preferred embodiment of the method for producing a high-strength cold-rolled steel sheet of the present invention (hereinafter also simply referred to as “the production method of the present invention”) will be described.
The production method of the present invention generally includes the above-described high-strength cold-rolled steel sheet according to the present invention by sequentially subjecting a steel material having the above composition to hot rolling, pickling, cold rolling, and annealing. Is the way to get. And in the manufacturing method of this invention, the process of annealing is divided into two processes.

<Steel material>
The steel material is not particularly limited as long as it is a steel material having the above composition.
The melting method of the steel material is not particularly limited, and a known melting method using a converter or an electric furnace can be employed. From the viewpoint of productivity and the like, it is preferable to make a slab (steel material) by a continuous casting method after melting, but it can also be made into a slab by a known casting method such as an ingot-bundling rolling method or a continuous slab casting method. Good.

<Hot rolling process>
A hot rolling process is a process of obtaining a hot-rolled sheet by hot-rolling the steel raw material which has the said composition.
The hot rolling process is not particularly limited as long as it is a process in which a steel material having the above composition is heated and subjected to hot rolling to obtain a hot rolled sheet having a predetermined size, and a normal hot rolling process is applied. it can.
As a normal hot rolling process, for example, a steel material is heated to a heating temperature of 1100 ° C. or more and 1300 ° C. or less, and hot rolling is performed on the heated steel material at a finish rolling outlet temperature of 850 ° C. or more and 950 ° C. or less. After the hot rolling is finished, cooling after appropriate rolling (specifically, for example, a temperature range of 450 ° C. or more and 950 ° C. or less is performed at an average cooling rate of 20 ° C./s or more and 100 ° C./s or less). An example is a hot rolling process in which cooling is performed after cooling and winding is performed at a coiling temperature of 400 ° C. or more and 700 ° C. or less to obtain a hot-rolled sheet having a predetermined size and shape.

<Pickling process>
The pickling step is a step of pickling the hot-rolled sheet obtained through the hot rolling step.
The pickling step is not particularly limited as long as it can be pickled to such an extent that cold rolling can be performed on the hot-rolled sheet. For example, a conventional pickling step using hydrochloric acid or sulfuric acid can be applied.

<Cold rolling process>
The cold rolling process is a process of performing cold rolling on the hot-rolled sheet that has undergone the pickling process. More specifically, the cold rolling step is a step of obtaining a cold rolled plate having a predetermined thickness by subjecting the hot rolled plate subjected to pickling to cold rolling with a rolling reduction of 30% or more.

<Cold rolling reduction: 30% or more>
The rolling reduction of cold rolling is 30% or more. When the rolling reduction is less than 30%, the processing amount is insufficient, and the number of austenite nucleation sites decreases. For this reason, austenite becomes coarse and non-uniform in the first-stage annealing process of the next process, and the lower bainite transformation in the holding process of the subsequent first-stage annealing process is suppressed, and martensite is generated excessively. As a result, the microstructure of the steel sheet after the first stage annealing process cannot be made a microstructure mainly composed of lower bainite. The portion that is martensite after the first stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second stage annealing step.
On the other hand, the upper limit of the rolling reduction is determined by the capability of the cold rolling mill, but if the rolling reduction is too high, the rolling load increases and the productivity may decrease. For this reason, the rolling reduction is preferably 70% or less.
The number of rolling passes and the rolling reduction per pass are not particularly limited.

<Annealing process>
An annealing process is a process which anneals the cold-rolled sheet obtained through the cold rolling process, and is a process including the 1st stage annealing process and 2nd stage annealing process mentioned later in detail.

<< First stage annealing process >>
In the first stage annealing step, the cold-rolled sheet obtained through the cold rolling step is heated at an annealing temperature T _{1 of} Ac ₃ points or more and 950 ° C. or less, and from the annealing temperature T ₁ , an average of more than 10 ° C./s at a cooling rate, cooling to cooling stop temperature T ₂ less than 250 ° C. or higher 350 ° C., by holding at the cooling stop temperature T ₂ 10s or more, a step of obtaining a first Danhiyanobe annealed sheets.
The purpose of this process is to make the microstructure of the steel sheet at the completion of the first stage annealing process into lower bainite. In particular, the portion that is martensite after the first-stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second-stage annealing step, so that the martensite is excessive in the first-stage annealing step. When it produces | generates, it will become difficult to obtain the microstructure of a desired steel plate.
By controlling the production conditions within the above range, a steel sheet having a microstructure mainly composed of lower bainite can be obtained, and the microstructure of the steel sheet after the second stage annealing step can be changed to a desired microstructure.

(Ac ₃ points)
Ac ₃ points (unit: ° C.) can be obtained from the following formula of Andrews et al.
^{_{Ac 3 = 910-203 [C] 1/2}} +45 [Si] -30 [Mn] -20 [Cu] -15 [Ni] +11 [Cr] +32 [Mo] +104 [V] +400 [Ti] +460 [Al ]
The parentheses in the above formula represent the content (unit: mass%) of the element in the parentheses in the steel sheet. When no element is contained, it is calculated as 0.

(Annealing temperature T ₁ : Ac ₃ points or more and 950 ° C. or less)
When the annealing temperature T ₁ is is Ac less than ₃ points, ferrite will remain in the annealing, ferrite ferrite remaining in the annealing in the subsequent cooling process the nucleus will grow. Thereby, since C distributes in austenite, lower bainite transformation is suppressed in the subsequent holding process, martensite is excessively generated, and the microstructure of the steel sheet after the first stage annealing process is mainly composed of lower bainite. It cannot be made into a microstructure.
On the other hand, the annealing temperature T ₁ is excessively coarsened austenite grains exceeds 950 ° C., since the formation of lower bainite is suppressed in the course retained after cooling, because the martensite excessively generated, the first stage annealing step The microstructure of the later steel sheet cannot be made into a microstructure mainly composed of lower bainite.
The portion that is martensite after the first stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second stage annealing step.
Thus, annealing temperatures _{T 1} is Ac ₃ point or more 950 ° C. or less.

Holding time at the annealing temperatures _{T 1} is not particularly limited, for example, is 10s or 1000s or less.

(Average cooling rate from the annealing temperature _{T 1} of to the cooling stop temperature _{T 2:} 10 ℃ / s greater)
If the average cooling rate from the annealing temperature T ₁ of to the cooling stop temperature T ₂ is less than 10 ° C. / s, ferrite is formed during cooling. Thereby, since C distributes in austenite, lower bainite transformation is suppressed in the subsequent holding process, martensite is excessively generated, and the microstructure of the steel sheet after the first stage annealing process is mainly composed of lower bainite. It cannot be made into a microstructure. The portion that is martensite after the first stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second stage annealing step.
Therefore, the average cooling rate from the annealing temperature T ₁ of to the cooling stop temperature T ₂ is 10 ° C. / s greater, preferably 15 ° C. / s or higher.
The upper limit of the average cooling rate is not particularly limited, but an excessively large cooling device is required to ensure an excessively high cooling rate. From the viewpoint of production technology and capital investment, the average cooling rate is 50 It is preferably at most ° C / s.
The cooling is preferably gas cooling, but can be performed by combining furnace cooling and mist cooling.

(Cooling stop temperature T ₂ : 250 ° C. or higher and lower than 350 ° C.)
Cooling the stop temperature T ₂ is less than 250 ° C., martensite microstructure of the steel sheet is excessively formed. The portion that is martensite after the first stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second stage annealing step.
On the other hand, at the cooling stop temperature T ₂ is 350 ° C. or more, the upper bainite is generated instead of lower bainite. The upper bainite is significantly coarser in the same Bain group size than the lower bainite. Therefore, a large number of retained austenite having an aspect ratio of 0.6 or less is generated inside the same Bain group after the subsequent second stage annealing step. The microstructure of the steel sheet after the second stage annealing step does not become the desired microstructure.
Therefore, the cooling stop temperature _{T 2} is less than 250 ° C. or higher 350 ° C.. More preferably, it is 270 degreeC or more and 340 degrees C or less.

(Retention time in the cooling stop temperature _{T 2:} 10s or more)
Is less than the holding time at the cooling stop temperature T ₂ is 10s (seconds), no lower bainite transformation is completed sufficiently. For this reason, martensite is excessively generated, and a desired microstructure cannot be obtained in the subsequent second stage annealing step. The portion that is martensite after the first stage annealing step tends to generate retained austenite having an aspect ratio of more than 0.6 in the subsequent second stage annealing step.
Therefore, the holding time at the cooling stop temperature _{T 2} is 10s or more. Preferably it is 30 s or more.
The upper limit of the holding time at the cooling stop temperature T ₂ is not particularly limited, if it is excessively held for a long time, as well as it requires a long production facilities, because the productivity of the steel sheet is remarkably reduced, 1800 s or less Is preferred.

After holding in the cooling stop temperature T _2, until the second stage annealing step following step, for example it may be cooled to room temperature, subsequently it is subjected to heating and second stage annealing step without cooling. In order to perform continuously without cooling to room temperature between the first stage annealing process and the second stage annealing process, two heating furnaces of a normal continuous annealing facility (CAL) are required in one line. After performing the first stage annealing process by CAL, the second stage annealing process is performed by passing the CAL once again.

<< Second stage annealing process >>
The second stage annealing process, a first Danhiyanobe annealed sheets obtained through the first-stage annealing process, was heated at 680 ° C. or higher 820 ° C. below the annealing temperature T ₃ (reheat), from annealing temperature T ₃ , by cooling to 300 of the cooling stop ° C. or higher 500 ° C. or less temperature T _4, which is a step of obtaining a second Danhiyanobe annealed sheets.

(Annealing temperature _T 3: 680 ℃ more than 820 ℃ or less)
When the annealing temperature T ₃ is lower than 680 ° C., since no generating a sufficient amount of austenite during annealing, can not be secured retained austenite desired amount the microstructure of the steel sheet after the second stage annealing step, the ferrite is excessive Become.
On the other hand, if the annealing temperature T ₃ is higher than 820 ° C., austenite excessively generated, effects of the second stage annealing before microstructure control from being initialized. For this reason, the ratio of the retained austenite having an aspect ratio of 0.6 or less and the ratio of the remaining austenite having an aspect ratio of 0.6 or less at the Bain group boundary may be set to a desired value. It becomes difficult.
Therefore, the annealing temperature _{T 3} is at 680 ° C. or higher 820 ° C. or less, preferably 700 ° C. or higher 800 ° C. or less.

Holding time at the annealing temperature _{T 3} is not particularly limited, for example, is 10s or 1000s or less.
The average cooling rate from the annealing temperature T ₃ to a cooling stop temperature T ₄ is not particularly limited, for example, is 50 ° C. / s or less 5 ° C. / s or higher.

(Cooling stop temperature T ₄ : 300 ° C. or more and 500 ° C. or less)
If the cooling stop temperature T ₄ is lower than 300 ° C., enrichment of C into austenite becomes insufficient, a large amount of martensite with retained austenite amount decreases is produced, it can not be obtained microstructure of the desired steel sheet .
On the other hand, the cooling stop temperature T ₄ is more than 500 ° C., since pearlite from austenite is generated, the amount of retained austenite is reduced, can not be obtained microstructure of the desired steel sheet.

Holding time at the cooling stop temperature _{T 4} is not particularly limited, for example, is 10s or 1800s or less.

Second Danhiyanobe annealed sheet after holding in the cooling stop temperature T ₄ is preferably cooled. This cooling is not particularly limited, and the cooling can be performed to a desired temperature such as room temperature by an arbitrary method such as cooling.

  When the plating process described later is not performed, the second-stage cold-rolled annealed sheet obtained through the second-stage annealing process becomes the high-strength cold-rolled steel sheet of the present invention.

<Plating process>
The second-stage cold-rolled annealed plate obtained through the second-stage annealing step may be further subjected to a plating treatment to form a plating layer on the surface thereof. In this case, the second-stage cold-rolled annealed plate having a plating layer formed on the surface is the high-strength cold-rolled steel plate of the present invention.

As the plating treatment, hot dip galvanizing treatment, hot dip galvanizing treatment and alloying treatment, or electrogalvanizing treatment is preferable. The hot dip galvanizing treatment, the hot dip galvanizing treatment and the alloying treatment, and the electrogalvanizing treatment are not particularly limited, and are conventionally known hot dip galvanizing treatment, conventionally known hot dip galvanizing treatment and alloying treatment, respectively. In addition, a conventionally known electrogalvanizing treatment is preferably used.
Prior to the plating treatment, pretreatment such as degreasing and phosphate treatment may be performed.

As the hot dip galvanizing treatment, for example, a conventional continuous hot dip galvanizing line is used to immerse the second stage cold-rolled annealing plate in a hot dip galvanizing bath and form a predetermined amount of hot dip galvanized layer on the surface. It is preferable that
When immersed in a hot dip galvanizing bath, the temperature of the second stage cold-rolled annealed plate is not less than the temperature of the hot dip galvanizing bath temperature −50 ° C. and not more than the temperature of the hot dip galvanizing bath temperature + 80 ° C. by reheating or cooling. It is preferable to adjust within the range.
The temperature of the hot dip galvanizing bath is preferably 440 ° C. or higher and 500 ° C. or lower.
The hot dip galvanizing bath may contain Al, Fe, Mg, Si or the like in addition to pure zinc.
The adhesion amount of the hot-dip galvanized layer can be adjusted to a desired adhesion amount by adjusting gas wiping or the like, and is preferably about 45 g / m ² per side.

The plated layer (hot galvanized layer) formed by the hot dip galvanizing process may be an alloyed hot dip galvanized layer by performing a usual alloying process as necessary.
The temperature for the alloying treatment is preferably 460 ° C. or more and 600 ° C. or less.
In the case of an alloyed hot dip galvanized layer, adjusting the effective Al concentration in the hot dip galvanizing bath to a range of 0.10% by mass or more and 0.22% by mass or less from the viewpoint of securing a desired plating appearance. preferable.

The electrogalvanizing treatment is preferably, for example, a treatment of forming a predetermined amount of electrogalvanized layer on the surface of the second stage cold-rolled annealed plate using a conventional electrogalvanizing line.
The adhesion amount of the electrogalvanized layer can be adjusted to a predetermined adhesion amount by adjusting the sheet passing speed or the current value, and is preferably about 30 g / m ² per side.

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

<Manufacture of cold-rolled steel sheet>
Molten steel having the composition shown in Table 1 below was melted by a generally known method and continuously cast to obtain a slab (steel material) having a thickness of 300 mm. A hot-rolled sheet was obtained by subjecting the obtained slab to hot rolling. The obtained hot-rolled sheet is pickled by a generally known method, and then cold-rolled at the rolling reduction shown in Tables 2 to 3 below to obtain a cold-rolled sheet (sheet thickness: 1.4 mm). It was.

The obtained cold-rolled sheet was annealed under the conditions shown in Tables 2 to 3 below to obtain a second-stage cold-rolled annealed sheet.
The annealing process was a two-stage process consisting of a first stage annealing process and a second stage annealing process.
It cooled to room temperature between the 1st stage annealing process and the 2nd stage annealing process.
Holding time at the annealing temperature T ₁ of the first stage annealing process was 100s.
In the second stage annealing step, the holding time at the annealing temperature T ₃ is 100 s, the average cooling rate from the annealing temperature T ₃ to the cooling stop temperature T ₄ is 20 ° C./s, and the holding time at the cooling stop temperature T ₄ Was 250 s.

Some of the second-stage cold-rolled annealed sheets were subjected to hot dip galvanizing treatment after the end of annealing, thereby forming a hot dip galvanized layer on the surface to obtain hot dip galvanized steel sheets.
In the hot dip galvanizing treatment, the second-stage cold-rolled annealed plate is reheated to a temperature in the range of 430 ° C. or higher and 480 ° C. or lower as necessary using a continuous hot dip galvanizing line, and a hot dip galvanizing bath (bath temperature) is used. : 470 ° C.) and adjusted so that the adhesion amount of the plating layer was 45 g / m ² per side. The bath composition was Zn-0.18 mass% Al.
At this time, in some hot-dip galvanized steel sheets, the bath composition was Zn-0.14 mass% Al, and after the plating process, alloying was performed at 520 ° C. to obtain alloyed hot-dip galvanized steel sheets.
The Fe concentration in the plating layer was 9% by mass or more and 12% by mass or less.
Another part of the second-stage cold-rolled annealed plate is subjected to an electrogalvanizing treatment after the annealing, and further using an electrogalvanizing line so that the amount of plating is 30 g / m ² per side. To give an electrogalvanized steel sheet.

In Tables 4 to 5 below, the second-stage cold-rolled annealed sheet that does not form a plating layer is “CR”, the hot-dip galvanized steel sheet is “GI”, the galvannealed steel sheet is “GA”, and the electrogalvanized steel sheet Was written as “EG”.
Hereinafter, a second-stage cold-rolled annealed plate that does not form a plated layer, and a second-stage cold-rolled annealed plate that forms a plated layer (hot-dip galvanized steel sheet, alloyed hot-dip galvanized steel sheet, and electrogalvanized steel sheet) These are collectively referred to as “cold rolled steel sheet”.
A cold-rolled steel sheet was produced as described above.

<Evaluation>
A test piece was collected from the obtained cold-rolled steel sheet and subjected to microstructure observation, measurement of the retained austenite area ratio, and tensile test. The test method was as follows.

<< Microstructure observation >>
First, a specimen for microstructural observation was collected from the cold rolled steel sheet.
Next, the collected specimen was polished so that the position corresponding to 1/4 of the plate thickness in the rolling direction cross section (L cross section) was the observation surface. The observation surface is corroded (1% by volume nital liquid corrosion), and then observed using a scanning electron microscope (SEM, magnification: 3000 times) in a visual field range of 30 μm × 35 μm, and imaged. SEM images were obtained.
The area ratio of each tissue was determined by image analysis using the obtained SEM image. The area ratio was an average value of 10 fields of view. In the SEM image, ferrite and bainitic ferrite are gray, and martensite and retained austenite are white. Therefore, each structure was judged from the color tone. Although it is difficult to accurately distinguish between ferrite and bainitic ferrite, the sum of these structures is important here, so the area ratio of the sum of ferrite and bainitic ferrite is not particularly distinguished. Asked.
The area ratio of retained austenite separately obtained by X-ray diffraction was subtracted from the area ratio of the white-colored structure to obtain the martensite area ratio. The volume ratio of austenite obtained by X-ray diffraction was treated as being equal to the area ratio.

  Furthermore, the test piece was polished by colloidal silica vibration polishing so that the position corresponding to 1/4 of the plate thickness in the rolling direction cross section (L cross section) became the observation surface. The observation surface was a mirror surface. Next, the processing transformation phase of the observation surface due to polishing strain was removed by ultra-low acceleration ion milling, and then electron beam backscatter diffraction (EBSD) measurement was performed to obtain local crystal orientation data. At this time, the SEM magnification was 1500 times, the step size was 0.04 μm, the measurement area was 40 μm square, and the WD was 15 mm. Analysis software: Using OIM Analysis 7, the obtained local orientation data was analyzed. The analysis was performed for three visual fields, and the average value was used.

Prior to data analysis, Grain Dilution function (Grain Tolerance Angle: 5 °, Minimum Grain Size: 5, Single Iteration: ON) and Grain CI Standardization Function (Grain Tolerance Angle: 5 °, Grain Tolerance Angle: 5 °, Grain Tolerance Angle: 5 °) The clean-up process according to) was performed once in order. Thereafter, only measurement points with CI values> 0.1 were used for analysis.
The fcc phase data was analyzed using the Area Fraction of the Grain Shape Aspect Ratio chart, and the ratio (area ratio) of retained austenite having an aspect ratio of 0.5 or less was determined from the retained austenite. In the above analysis, Method 2 was used as the grain shape calculation method.
Further, for the data of the bcc phase, after highlighting the region belonging to the same Bain group with the same color using the highlight function, the different austenite with the aspect ratio obtained earlier of 0.5 or less is used in a different color. The ratio of those existing at the boundary of the colored region, that is, the Bain group boundary (including the former austenite grain boundary) was obtained as an area ratio.

<Measurement of residual austenite area ratio>
A specimen for X-ray diffraction is taken from a cold-rolled steel sheet, ground and polished so that the position corresponding to 1/4 of the plate thickness becomes the measurement surface, and diffracted X-ray intensity is measured by the X-ray diffraction method. From the volume fraction of retained austenite. CoKα rays were used as incident X-rays.
When calculating the volume fraction of retained austenite, the {111}, {200}, {220} and {311} faces of the fcc phase (residual austenite) and the {110}, {200} and {211} of the bcc phase The intensity ratio was calculated for all combinations of the integrated intensity of the peak of the surface, the average value thereof was obtained, and the volume ratio of retained austenite was calculated.
The volume ratio of austenite thus obtained was defined as the area ratio.

<Tensile test>
A JIS No. 5 tensile test piece (JIS Z 2001) having a tensile direction in the direction perpendicular to the rolling direction (C direction) is taken from the cold-rolled steel sheet and subjected to a tensile test in accordance with the provisions of JIS Z 2241. Strength (TS) and elongation (El) were measured. Ten tensile tests were performed in total, and the average value of the measurement results was taken as the values of TS and El.

(Strength)
The case where TS was 1320 MPa or more was evaluated as high strength.

(Ductility)
When TS is 1320 MPa or more and less than 1460 MPa, El is 15% or more, and when TS is 1460 MPa or more, El is 13% or more, it was evaluated as high ductility (good ductility).

(Ductility anisotropy)
Further, from a cold-rolled steel sheet, a JIS No. 5 tensile test piece (JIS Z 2001) having a direction parallel to the rolling direction (L direction) as a tensile direction and a direction at 45 ° (D direction) with respect to the rolling direction are pulled. A JIS No. 5 tensile test piece (JIS Z 2001) having a direction was taken, a tensile test based on the JIS Z 2241 rule was performed, and the elongation (El) was measured. Ten tensile tests were performed in total, and the average value of the measurement results was taken as the value of El.

From the obtained elongation (El), ΔEl defined by the following formula (X) was calculated, and the anisotropy of the elongation was evaluated.
ΔEl = | (El _L + El _C −2El _D ) / 2 | (X)
(In the formula (X), ΔEl: elongation anisotropy (unit:%), El _L : elongation in the L direction (unit:%), El _C : elongation in the C direction (unit:%), El _D : D Directional elongation (unit:%))
When TS is 1320 MPa or more and less than 1460 MPa, when ΔEl is 7% or less, and when TS is 1460 MPa or more and ΔEl is 5% or less, it was evaluated that ductility anisotropy was small.

FIG. 2 is a graph in which a part of the results in Tables 4 to 5 is plotted. More specifically, FIG. 2 shows the proportion of residual austenite with an aspect ratio of 0.6 or less, and the proportion of residual austenite with an aspect ratio of 0.6 or less that exists at the Bain group boundary. It is a graph which shows the influence which has on the anisotropy of elongation.
As can be seen from the graph of FIG. 2, the ratio of residual austenite having an aspect ratio of 0.6 or less is 70% or more, and the residual austenite having an aspect ratio of 0.6 or less exists at the Bain group boundary. Only when the ratio of the material is 50% or more, a steel sheet having low elongation anisotropy is obtained.

As is apparent from Tables 1 to 5 and FIG. 2, the cold-rolled steel sheets of the examples of the present invention all have high strength with a tensile strength (TS) of 1320 MPa or more, and also have high ductility. Furthermore, the ductility anisotropy is small.
On the other hand, in the comparative example, any of the above characteristics was insufficient.

B1, B2, B3: Bain group RA _1, RA _2: residual austenite

Claims (5)

% By mass
C: more than 0.18% and 0.45% or less,
Si: 0.50% or more and 2.50% or less,
Mn: more than 2.50% and 4.00% or less,
P: 0.050% or less,
S: 0.0100% or less,
Al: 0.010% or more and 0.100% or less, and
N: including 0.0100% or less, having a composition consisting of the balance Fe and inevitable impurities,
C and Mn of the composition satisfy the following formula (1) in mass%,
In the microstructure, the sum of the area ratios of ferrite and bainitic ferrite is 10% or more and 50% or less, the area ratio of retained austenite is more than 15% and 50% or less, and the area ratio of martensite is 15%. 60% or less,
Of the retained austenite, the proportion of those having an aspect ratio of 0.6 or less is 70% or more in area ratio,
A high-strength cold-rolled steel sheet having an area ratio of 50% or more of residual austenite having an aspect ratio of 0.6 or less and existing at the boundary of the Bain group.
7.5 × C + Mn ≧ 5.0 (1)
However, in Formula (1), C and Mn show content of each element. The composition is further in mass%,
Ti: 0.005% or more and 0.035% or less,
Nb: 0.005% or more and 0.035% or less,
V: 0.005% or more and 0.035% or less,
Mo: 0.005% or more and 0.035% or less,
B: 0.0003% or more and 0.0100% or less,
Cr: 0.05% or more and 1.00% or less,
Ni: 0.05% or more and 1.00% or less,
Cu: 0.05% or more and 1.00% or less,
Sb: 0.002% or more and 0.050% or less,
Sn: 0.002% to 0.050%,
Ca: 0.0005% or more and 0.0050% or less,
Mg: 0.0005% or more and 0.0050% or less, and
The high-strength cold-rolled steel sheet according to claim 1, comprising at least one element selected from the group consisting of REM: 0.0005% to 0.0050%.   The high-strength cold-rolled steel sheet according to claim 1 or 2, wherein the surface has a plating layer. A method for producing the high-strength cold-rolled steel sheet according to any one of claims 1 to 3,
A hot rolling process for obtaining a hot-rolled sheet by subjecting the steel material having the composition according to claim 1 or 2 to hot rolling,
Pickling step of pickling the hot-rolled sheet;
A cold rolling step of obtaining a cold-rolled sheet by subjecting the hot-rolled sheet subjected to the pickling to cold rolling with a rolling reduction of 30% or more;
The cold-rolled sheet is heated at an annealing temperature T ₁ of Ac ₃ points or more and 950 ° C. or less, and a cooling stop temperature of 250 ° C. or more and less than 350 ° C. at an average cooling rate of more than 10 ° C./s from the annealing temperature T _1. It cooled to T _2, by holding in the cooling stop temperature T ₂ 10s or more, and the first stage annealing step of obtaining a first Danhiyanobe annealed sheets,
Said first Danhiyanobe annealed sheet, by heating at an annealing temperature _{T 3} of 680 ° C. or higher 820 ° C. or less, cooling from the annealing temperature _{T 3,} to 300 ° C. or higher 500 ° C. or less of the cooling stop temperature _{T 4,} A second-stage annealing step to obtain a second-stage cold-rolled annealing plate;
A method for producing a high-strength cold-rolled steel sheet.   5. The high-strength cold-rolled steel sheet according to claim 4, further comprising a plating step of subjecting the second-stage cold-rolled annealed plate to a hot-dip galvanizing treatment, a hot-dip galvanizing treatment and an alloying treatment, or an electrogalvanizing treatment. Production method.