JP2011179071A - High-tensile cold-rolled steel sheet and method of producing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-tensile cold-rolled steel sheet having &ge;500 MPa tensile strength and excellent in elongation, stretch-flangeability and balance between the strength, extension and stretch-flangeability. <P>SOLUTION: The high-tensile cold-rolled steel sheet is produced by subjecting a steel having composition at least containing, by mass, 0.03 to 0.20% C, 0.01 to 1.5% Si, &le;1.0% Mn, &le;0.08% P, &le;0.005% S, 0.01 to 0.08% Al, 0.001 to 0.005% N, one or more selected from Ti, Nb and V in an amount of 0.02 to 1.0% in total, and the balance Fe with inevitable impurities to an annealing treatment of a prescribed condition, after hot-rolling and cold-rolling in order. The produced high-tensile cold-rolled steel sheet has a composite structure composed of two ferritic phases with widely different strengths, wherein the soft ferritic phase with a lower strength has a grain diameter of &le;10 &mu;m, and &ge;60% crystal grains of the soft ferritic grains are not in contact with the crystal grains of the other soft ferritic phase. <P>COPYRIGHT: (C)2011,JPO&amp;INPIT

Description

The present invention relates to a high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more, which is suitable mainly for use in automobile body parts and the like, and in particular, high tension excellent in stretch flangeability and strength-stretch-stretch flangeability balance. The present invention relates to a cold-rolled steel sheet and a manufacturing method thereof.
“Excellent stretch flangeability” means a case where the hole expansion ratio λ is 100% or more, and “excellent stretch property” means a case where the elongation El is 30% or more. “Excellent strength-elongation-stretch flangeability balance” means that the product TS × El × λ of the tensile strength TS, total elongation El, and hole expansion ratio λ is 2000000 MPa% ² or more.

In recent years, in order to suppress the discharge amount of carbon dioxide from automobiles, the weight reduction of automobile bodies has been promoted using high-tensile steel plates. Further, in order to ensure the safety of passengers, high-strength steel plates having a TS of about 590 to 780 MPa are often used for automobile bodies. In the future, it is expected that steel sheets with a strength of 900 MPa or more will be used more frequently as the strength further increases.
However, since many automotive body parts made of steel sheets are formed by press working, high-tensile steel sheets used for body parts are required to have excellent press formability. Therefore, the mechanical properties of the steel sheet are required to have high stretch flangeability (hole expansion ratio λ) and high ductility while having high strength TS.

As a steel sheet having high strength and high ductility, a ferrite-martensite composite structure steel sheet (Dual-Phase (DP) steel sheet) in which a parent phase is a ferrite structure and martensite is dispersed in the ferrite structure is known ( For example, see Patent Document 1). This DP steel sheet has a high elongation El due to the ferrite phase which is a soft layer while realizing high strength by containing martensite which is a hard phase.
However, DP steel sheet has a large difference in deformability. As a result of the mixture of martensite phase and ferrite phase, void formation and crack propagation at the interface between martensite and ferrite are easy. There is a problem that is bad.

Therefore, in order to improve the stretch flangeability of DP steel, steel sheets have been developed in which the hardness of martensite in DP steel is reduced by adding temper annealing to DP steel, and the hardness difference between ferrite and martensite is reduced. (For example, Patent Documents 2 and 3).
However, even when tempering is performed to reduce the hardness of martensite, the deformability of martensite is not greatly improved, and the difference in deformability from the ferrite phase is still large. Sex is bad. Moreover, since the number of tempering annealing steps increases in addition to the normal steps, it is disadvantageous in terms of cost.

Recently, TRIP steel sheets have attracted attention. The TRIP steel sheet exhibits excellent ductility by generating retained austenite in a ferrite structure or a multiphase structure of ferrite, bainite, and martensite, and this retained austenite undergoes strain-induced transformation during work deformation. For example, in Patent Document 4, a material having an excellent strength-elongation balance of TS: 108 MPa and El: 22% is obtained.
However, this TRIP steel sheet also has a defect that it is inferior in stretch flangeability because fracture tends to proceed at the interface between martensite and matrix structure generated by strain-induced transformation (in the example of Patent Document 4, TS When λ is 108 MPa, λ: 20%), application applications are limited.

  In view of this, various studies have been made to provide a steel sheet having excellent formability such as stretch flangeability while maintaining an excellent balance between strength and elongation due to retained austenite. For example, as described in Patent Document 5, a TRIP steel sheet having tempered martensite and tempered bainite as a parent phase structure and retained austenite as a second phase structure is disclosed. However, even in these steels, the hole expansion ratio λ is about 50%, which is insufficient for pressing under severe conditions.

In addition, by dispersing a fine residual austenite phase with a size of several hundreds of nanometers, the size of martensite generated by strain-induced transformation is reduced, and the main purpose is to suppress destruction in the vicinity of martensite. A steel plate has been invented (Patent Document 6).
However, in order to produce a steel containing such finely dispersed residual austenite phase, (1) since expensive elements such as Co, Ni, Ag, and Pt are added as austenite stabilizing elements, the cost increases. (2) Solution treatment at 1270 ° C. for 5 hours or longer and annealing for a long time to segregate the austenite stabilizing elements are necessary, and the annealing time needs to be strictly controlled, and the process is too complicated. (3) It is necessary to add Si in order to secure retained austenite and cannot be applied to the plated steel sheet. (4) Residual austenite is too small to cause strain-induced transformation. There are problems such as high elongation, which is a characteristic of steel, difficult to develop.

  On the other hand, as described in Patent Document 7, as a steel sheet having high strength and high stretch flangeability, there is a bainite steel sheet (λ: 75% when TS: 755 MPa). However, since the bainite structure is oriented to improve the stretch flangeability, the elongation value is low (El: 23% when TS: 755 MPa), and the application is limited in reality. is there.

JP 55-122821 JP-A-5-311244 JP 2004-52071 A JP-A-9-104947 JP 2002-309334 A JP 2005-179703 A Japanese Patent Laid-Open No. 3-180426

  An object of the present invention is a high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more which is suitable mainly for the use of automobile body parts and the like, and is excellent in elongation, stretch flangeability, strength-stretch-stretch flangeability balance It is to provide a tension cold-rolled steel sheet and a manufacturing method thereof.

The gist of the present invention is as follows.
(1) High-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more, and in mass%, C: 0.03 to 0.20%, Si: 0.01 to 1.5%, Mn: 1.0 %: P: 0.08% or less, S: 0.005% or less, Al: 0.01-0.08%, N: 0.001-0.005%, Ti, Nb, V A composition comprising at least 0.02 to 1.0% of a total of two or more species, the balance being composed of Fe and unavoidable impurities, and a composite structure consisting of two types of ferrite phases that differ greatly in strength, The grain size of the soft ferrite phase, which is a low strength ferrite phase, is 10 μm or less, and 60% or more of the crystal grains of the soft ferrite phase are not in contact with other soft ferrite phase grains. High tensile cold-rolled steel sheet.
(2) The high-strength cold rolling according to claim 1, wherein the composition further includes 0.1% or less in total of one or more of Ca and REM in mass%. steel sheet.

(3) A method for producing a high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more, wherein C: 0.03 to 0.20%, Si: 0.01 to 1.5%, Mn: 1.0% or less, P: 0.08% or less, S: 0.005% or less, Al: 0.01 to 0.08%, N: 0.001 to 0.005%, Ti, Nb, V A steel slab having a composition containing at least 0.02 to 1.0% of one or more of them in total, the balance being Fe and inevitable impurities, after heating to a heating temperature of 1100 ° C. or higher, A hot rolling step of rough rolling into a sheet bar, subjecting the sheet bar to a finish rolling exit temperature: 900 ° C. or higher, and a coiling temperature: 750 ° C. or lower to be a wound hot rolled sheet, A cold rolling step of pickling and cold rolling the hot rolled sheet to form a cold rolled sheet, and the cold rolled sheet After heating to an annealing temperature in the temperature range of (Ac3 transformation point temperature) to (Ac3 transformation point temperature + 50 ° C.) and holding for 10 to 120 s, from the annealing temperature (Ar3 transformation point temperature −150 ° C.) to Cooling is performed so that the average cooling rate up to a predetermined temperature in the temperature range of (Ar3 transformation point temperature −30 ° C.) is 20 ° C./s or more, and the predetermined temperature is maintained for 500 to 1000 s. Thereafter, from the annealing temperature (Ar3 transformation) An annealing process in which cooling is performed so that an average cooling rate to a predetermined temperature in a temperature range of (point temperature −250 ° C.) to (Ar 3 transformation point temperature −150 ° C.) is 20 ° C./s or more, and held at the predetermined temperature for 10 to 1000 seconds The manufacturing method of the high tension cold-rolled steel sheet characterized by performing these.
(4) The high steel according to claim 3, wherein the steel slab has a composition that further includes 0.1% or less in total of one or more of Ca and REM in mass%. A method for producing a tension cold-rolled steel sheet.

  According to the present invention, a high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more and a high-tensile cold-rolled steel sheet excellent in elongation, stretch flangeability, and strength-stretch-stretch flangeability balance can be produced.

As a result of continuing the experimental study to further increase the elongation and stretch flangeability of the high-strength steel sheet, the present inventors have found that the conventional DP steel sheet and TRIP steel sheet have good ductility (elongation) but stretch flangeability (hole The reason why the spreadability is inferior to λ) is that the deformability of the soft phase (mainly ferrite phase) and the hard phase (martensite originally contained, and martensite generated by strain-induced transformation of retained austenite) contained in these steel sheets. Due to the difference, stress concentration occurs in the vicinity of the interface between the soft phase and the hard phase, voids and cracks are generated in the soft phase near the interface, cracks easily propagate in the soft phase near the interface, and the steel sheet becomes early. It is because it destroys. "
Based on this idea, it has been found that the stretch flangeability can be remarkably improved if the crystal grains of the soft phase are not coarse and are not connected and isolated. Furthermore, it has been found that if the deformability of the hard phase is large, the stretch flangeability can be remarkably improved.

Hereinafter, the high-strength steel sheet of the present invention thus made will be described in detail.
First, the structure of the high-strength steel sheet of the present invention will be described.
In order to obtain good ductility (elongation), a soft phase having a large deformability is required. On the other hand, in order to obtain a high strength of 500 MPa or more, a hard phase is necessary. Furthermore, in order to obtain good stretch flangeability, the hard phase also needs a certain degree of deformability.
In the present invention, two types of ferrite phases (hereinafter sometimes referred to as a soft ferrite phase and a hard ferrite phase, respectively) having greatly different strengths are used as the hard layer and the soft phase, and the structure is composed of these ferrite layers. A complex organization is assumed.
As the soft ferrite phase, a ferrite phase having a low precipitate density was used, and as the hard ferrite phase, precipitation strengthened ferrite containing a high density precipitate was used. Precipitation strengthened ferrite has high strength due to precipitates dispersed at a high density and has a low dislocation density, and therefore has a greater deformability than a structure containing dislocations at a high density such as martensite.

The particle diameter of the soft ferrite phase is desirably 10 μm or less. When the steel sheet is stretched and deformed by flange, a crack generated near the interface between the soft phase and the hard phase easily propagates in the soft phase. That is, the size of the initial crack (initial crack) depends on the soft phase particle size. As the deformation progresses, the initial crack grows and eventually breaks. However, if the initial crack size is 10 μm or less, the crack is not easily broken. On the contrary, if the size of the initial crack is 10 μm or more, the initial crack easily develops and breaks.
Therefore, the particle diameter of the soft ferrite phase is desirably 10 μm or less. Note that the size of the crack that easily starts to grow also depends on the strength level of the steel sheet. Within the scope of the present invention, 10 μm is the critical size.

It is desirable that 60% or more of the soft ferrite phase crystal grains are not in contact with other soft ferrite phase crystal grains. If the ferrite grains are adjacent to each other, cracks generated in the vicinity of the interface with the hard phase can propagate in the range in which the soft ferrite grains continue, so that the cracks become large, leading to early breakage and poor stretch flangeability.
If soft ferrite grains are isolated in the hard phase, cracks that occur near the interface with the hard phase can only grow easily within the range of one ferrite grain, so the crack does not become large, and the fracture is delayed, Good stretch flangeability can be obtained. If 60% or more of all soft ferrite grains are in an isolated state, desired stretch flangeability can be obtained.
In conventional DP steel and TRIP steel, the soft phase is the parent phase, and the soft phase is connected to a very wide region (may be connected to the entire steel plate). For this reason, cracks are easily increased, leading to early breakage, and stretch flangeability is deteriorated.

Next, the chemical composition of the steel of the present invention will be described. Hereinafter, all the units of chemical components are mass%.
C: 0.03-0.20%
C is essential for bonding with alloy elements such as Ti, Nb, and V to precipitate fine carbides. In order to achieve a steel plate strength of 500 MPa or more, at least 0.03% is necessary. On the other hand, if C is excessive, defects due to center segregation are likely to occur at the casting stage, and weldability also deteriorates. Therefore, the upper limit was made 0.20%.
Si: 0.01 to 1.5%
Si is an element useful as a solid solution strengthening element of the soft ferrite phase. When the strength of the soft ferrite phase is increased, the strength difference from the hard phase is reduced, so that stress concentration is relaxed and stretch flangeability is improved. Therefore, 1.0% or more is necessary. However, when plating is performed as a plated steel sheet, the addition of Si significantly deteriorates the plating properties, so that it is preferable that the amount of Si is small. If it is about 0.01%, the plating property is not greatly affected. Therefore, when the plating property needs to be considered, the lower limit is set to 0.01%. On the other hand, if added over 1.5%, not only the effect of Si is saturated but also the workability deteriorates, so the upper limit was made 1.5%.

Mn: 1.0% or less Mn is useful as a solid solution strengthening element like Si. However, since the effect of suppressing the precipitation of fine carbides is strong, it cannot be added in a large amount in the present invention. If added over 1.0%, the amount of strengthening due to the precipitation of fine carbides becomes remarkably small, so the upper limit was made 1.0%.
P: 0.08% or less Since P deteriorates the workability at the time of hot rolling, the lower one is desirable. Since 0.08% or less has little influence, 0.08% was made the upper limit. More preferably, it is 0.03% or less.

S: 0.005% or less S forms Mn sulfide, which becomes a starting point of fracture and deteriorates stretch flangeability. Therefore, it is desirable that S is low. As will be described later, when Ca or REM is added, the formation of Mn sulfide is suppressed. Therefore, when an appropriate amount of Ca or REM is added, S is remarkable in stretch flangeability even if it is added to a maximum of 0.005%. There is no impact. Therefore, the upper limit of S is set to 0.005%. Desirably, it is 0.002% or less, and more desirably 0.001% or less.
Al: 0.01 to 0.08%
Al is added in an amount of 0.01% or more for deoxidation, but when the addition amount increases, inclusions such as alumina increase and stretch flangeability deteriorates, so 0.08% is made the upper limit.

N: 0.001 to 0.005%
Since N combines with Ti and precipitates as TiN, the effective amount of Ti that can be used for precipitation strengthening decreases, and if coarse TiN is present, it becomes a fracture starting point and deteriorates stretch flangeability. If it exceeds 0.005%, the stretch flangeability deteriorates remarkably, so 0.005% is made the upper limit. Moreover, since extremely low is economically disadvantageous, the lower limit is made 0.0001%.
Ti, Nb, V: 0.02 to 1.0% in total of 1 type or 2 types or more
Ti, Nb, and V need to be added in an amount of 0.02% or more for the purpose of strengthening the ferrite phase by carbide precipitation. However, when the addition amount increases, coarse inclusions such as TiN increase and stretch flangeability deteriorates, so 1.0% is made the upper limit.

Although the present invention is based on steel containing at least the above elements, the following elements can be further added.
Ca, REM: 1 type or 2 types or more in total 0.1% or less When Ca and REM are added, the production | generation of Mn sulfide is suppressed and stretch flangeability can be improved. This is particularly effective when the amount of S is large. However, if it is added in a large amount, inclusions such as Ca oxide are generated and serve as a starting point for fracture, which adversely affects stretch flangeability. Therefore, the upper limit is made 0.1% in total.

Next, a method for producing the steel of the present invention will be described.
The most important points to be noted in the production of the steel of the present invention are to produce two types of ferrite phases, soft and hard, and to keep the crystal grains of the soft ferrite phase as isolated as possible. Therefore, an important point as a manufacturing method is a process from annealing to cooling after cold rolling. The process from obtaining a cold-rolled sheet from a slab is the same as a normal cold-rolled sheet manufacturing process.
Below, each process is demonstrated.

Hot rolling process:
A steel slab having a specified composition is heated to a heating temperature of 1100 ° C. or higher and then roughly rolled to form a sheet bar. If the heating temperature at this time is less than 1100 ° C., the temperature at the end of hot rolling is less than 900 ° C., and the structure at the end of hot rolling includes unrecrystallized grains extending in the rolling direction, and mechanical characteristics, particularly hole expansion. Anisotropy occurs in the characteristics, and as a result, the characteristics deteriorate. Therefore, the heating temperature at the time of hot rolling needs to be 1100 ° C. or higher.
For the above reasons, the finish rolling outlet temperature needs to be 900 ° C. or higher. Furthermore, if the coiling temperature exceeds 750 ° C., the time required for the hot rolled coil to cool to room temperature becomes too long, which is economically disadvantageous. Therefore, the winding temperature of the hot-rolled sheet is set to 750 ° C. or lower.

Cold rolling process:
It is the same as a normal cold rolling process, and cold rolling is performed at a required reduction rate according to the application.
Annealing process after cold rolling:
In order to make the steel structure an austenite single phase, the cold-rolled sheet obtained in the cold-rolling step is heated to an annealing temperature T1 in a temperature range of (Ac3 transformation point temperature) to (Ac3 transformation point temperature + 50 ° C.) for 10 to 120 s. Hold.
If the annealing temperature T1 is equal to or lower than the Ac3 transformation point temperature, an austenite single-phase structure is not formed, so the lower limit temperature is the Ac3 transformation point temperature. When the annealing temperature T1 is too high, austenite crystal grains become coarse, the martensite structure after cooling becomes coarse, and stretch flangeability deteriorates. Therefore, the upper limit of the annealing temperature T1 is set to (Ac3 transformation point temperature + 50 ° C.). If the holding time t1 at the annealing temperature is less than 10 s, the austenite transformation may not be completed. Conversely, holding for more than 120 s is economically wasteful because the austenite transformation is completed and the austenite crystal grains only become coarse. Therefore, it was set to 120 s or less.

Cooling process:
(1) Hard ferrite production process:
In order to ensure strength, it is necessary to generate a hard ferrite phase with a certain fraction. After the holding at the annealing temperature T1 is completed, a predetermined temperature (hereinafter, hard ferrite generation processing temperature T2) in a temperature range from the annealing temperature to (Ar3 transformation point temperature −150 ° C.) to (Ar3 transformation point temperature −30 ° C.). By cooling to 500 ° C. and holding at that temperature for 500 to 1000 s (t2), hard ferrite can be generated.
When the steel is held at a temperature lower than the Ar3 transformation point temperature, transformation from the austenite phase to the ferrite phase starts, but ferrite is produced in the temperature range of (Ar3 transformation point temperature -150 ° C) to (Ar3 transformation point temperature-30 ° C). When transformation is performed, so-called phase interface precipitation occurs in which carbides such as Ti, Nb, and V are precipitated in the vicinity of the interface between the formed ferrite phase and austenite matrix almost simultaneously with the ferrite transformation. Then, the ferrite transformation progresses and the ferrite-austenite interface moves and repeats the precipitation. As a result, fine carbides are distributed in a high density inside the formed ferrite phase, and a hard ferrite phase is obtained.

Phase interface precipitation occurs only when the interfacial transfer rate and precipitation rate are well balanced, and therefore occurs only within a certain temperature range. Even when the ferrite transformation temperature is lower than (Ar3 transformation point temperature−150 ° C.) or higher than (Ar3 transformation point temperature−30 ° C.), phase interface precipitation does not occur. Accordingly, the hard ferrite generation treatment temperature T2 must be in the temperature range of (Ar3 transformation point temperature−150 ° C.) to (Ar3 transformation point temperature−30 ° C.).
As the hard ferrite generation treatment temperature T2 is lower and the holding time t2 at that temperature is longer, the hard ferrite fraction tends to be higher.
The holding time t2 needs to be set depending on the hard ferrite generation processing temperature T2, but if it is less than 500 s, the hard ferrite fraction is too low and it is highly possible that a steel plate strength of 500 MPa or more cannot be achieved. Are likely to connect without isolation. On the other hand, if it is held for more than 1000 s, the hard ferrite fraction is too high, the elongation and stretch flangeability deteriorate, and the process time becomes longer, which is economically disadvantageous. Accordingly, the holding time t2 is desirably 500 to 1000 s.
The cooling rate from the annealing temperature T1 to the hard ferrite generation treatment temperature T2 is 20 ° C./s or more on average. When the cooling rate is slow, ferrite transformation starts during cooling and phase interface precipitation does not occur, so that a predetermined structure cannot be obtained.

(2) Soft ferrite production process:
After generating hard ferrite by holding at a hard ferrite generation processing temperature T2 for a predetermined time t2, it is rapidly cooled to a predetermined temperature for generating soft ferrite (hereinafter referred to as soft ferrite generation processing temperature T3), and the remainder is made soft ferrite. A soft ferrite phase can be obtained by reducing the density of carbide precipitates and dislocations contained in the ferrite. Suppression of carbide precipitation can be realized by lowering the ferrite transformation temperature and making Ti, Nb, and V difficult to diffuse. Conversely, if the ferrite transformation temperature is too low, dislocations are introduced into the ferrite and the ferrite becomes hard. In the composition range of the steel of the present invention, if the ferrite transformation temperature is higher than (Ar3 transformation point temperature -150 ° C), carbide precipitates in the ferrite, and if it is lower than (Ar3 transformation point temperature -250 ° C), the dislocation density is high. Become. Therefore, the soft ferrite generation treatment temperature is set to (Ar3 transformation point temperature−250 ° C.) to (Ar3 transformation point temperature−150 ° C.).

The holding time t3 at the soft ferrite generation processing temperature T3 must be such that the remainder becomes ferrite. If untransformed austenite remains, martensite is generated after cooling to room temperature after the step. Since martensite is hard and has poor deformability, when stretched flange deformation is applied, cracks are generated near the interface with the ferrite and easily break.
The holding time t3 should be adjusted according to the desired strength, steel composition, hard ferrite generation temperature T2, and soft ferrite generation temperature T3, but if it is 10 s or less, there is a high possibility that austenite remains. It is difficult to accurately control the holding time t3. On the other hand, holding for 1000 seconds or more is disadvantageous economically because the process time becomes long. Accordingly, the holding time t3 at the soft ferrite generation processing temperature T3 is set to 10 to 1000 s.
The cooling rate from the annealing temperature T2 to the soft ferrite generation treatment temperature T3 is 20 ° C./s or more on average. If the cooling rate is slow, hard ferrite is generated during cooling, the hard ferrite fraction becomes too high, and ductility may deteriorate.

  The following examples further illustrate the feasibility and effects of the present invention.

A steel plate having a composition as shown in Table 1 is heated to 1100 to 1250 ° C., and hot-rolling is completed at 900 to 950 ° C., cooled to 500 to 700 ° C., wound up, pickled, and then cold-rolled to 1 .2 mm thickness. Then, Ac3 and Ar3 transformation temperature was calculated | required by calculation from the component (mass%) of each steel according to the following formula.
Ac3 = 910−203 × (C%) ^0.5 −15.2 × Ni% + 44.7 × Si% + 104 × V% + 31.5 × Mo% −30 × Mn% −11 × Cr% −20 × Cu% + 700 × P% + 400 × Al% + 400 × Ti%
Ar3 = 910−273 × (C%) + 25 × Si% −74 × Mn% −56 × Ni% −16 × Cr% −9 × Mo% −5 × Cu%

These steel plates were annealed under the conditions shown in Table 2. All annealing treatments were performed under an Ar atmosphere. JIS No. 5 tensile test specimens were collected from these steel plates and measured for mechanical properties. In addition, a hole expansion test was performed in accordance with the Steel Federation standard, and the hole expansion rate was obtained.
To distinguish between soft ferrite and hard ferrite, hardness was measured using a micro Vickers tester, and grains having a particle size of 200 Hv or more were regarded as hard ferrite. Moreover, in order to measure the particle diameter of hard ferrite and soft ferrite and to measure the ratio of isolated soft ferrite, observation with an optical microscope was performed. For observation with an optical microscope, a steel sample etched with a nital solution (nitric acid 5% + ethanol 95%) after mirror polishing was used.

Table 3 shows the mechanical properties and structural observation results of these steels.
It can be seen that all the steels of the present invention have a tensile strength of 500 MPa or more, excellent stretch flangeability (hole expansion ratio λ is 100% or more), and excellent strength-stretch-stretch flangeability balance. Moreover, the steel plate manufactured in the scope of the claims of the present invention has the above-described structure as observed with an optical microscope. On the other hand, the comparative example which does not satisfy the scope of the present invention has a strength of less than 500 MPa, or is inferior in strength-elongation-stretch flangeability balance.

Claims (4)

A high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more, in mass%,
C: 0.03 to 0.20%
Si: 0.01 to 1.5%,
Mn: 1.0% or less,
P: 0.08% or less,
S: 0.005% or less,
Al: 0.01 to 0.08%,
N: 0.001 to 0.005%,
0.02 to 1.0% in total of one or more of Ti, Nb, and V,
And having a composite structure consisting of two types of ferrite phases, the composition of which the balance is composed of Fe and inevitable impurities, and the strength is greatly different,
The grain size of the soft ferrite phase, which is a low strength ferrite phase, is 10 μm or less, and 60% or more of the crystal grains of the soft ferrite phase are not in contact with other soft ferrite phase grains. High tensile cold-rolled steel sheet.   2. The high-tensile cold-rolled steel sheet according to claim 1, wherein the composition further includes 0.1% or less in total of one or more of Ca and REM in mass%. A method for producing a high-tensile cold-rolled steel sheet having a tensile strength of 500 MPa or more, in mass%,
C: 0.03 to 0.20%
Si: 0.01 to 1.5%,
Mn: 1.0% or less, P: 0.08% or less,
S: 0.005% or less,
Al: 0.01 to 0.08%,
N: 0.001 to 0.005%,
0.02 to 1.0% in total of one or more of Ti, Nb, and V,
The steel slab having a composition consisting of at least Fe and the inevitable impurities is heated to a heating temperature of 1100 ° C. or higher, and then roughly rolled into a sheet bar. The finish rolling exit temperature of the sheet bar is 900 ° C. A hot rolling process in which finish rolling is performed as described above, and a coiling temperature is 750 ° C. or less to be a coiled hot rolled sheet,
A cold rolling step of pickling and cold rolling the hot rolled sheet to form a cold rolled sheet; and
The cold-rolled sheet is subjected to an annealing treatment in which it is heated to an annealing temperature in a temperature range of (Ac3 transformation point temperature) to (Ac3 transformation point temperature + 50 ° C.) and held for 10 to 120 seconds, and then from the annealing temperature (Ar3 transformation point temperature) −150 ° C.) to (Ar 3 transformation point temperature −30 ° C.), cooling is performed so that the average cooling rate to a predetermined temperature is 20 ° C./s or more, and the predetermined temperature is maintained for 500 to 1000 s, and then the annealing is performed. The cooling is performed so that the average cooling rate from the temperature to the predetermined temperature in the temperature range of (Ar3 transformation point temperature−250 ° C.) to (Ar3 transformation point temperature−150 ° C.) is 20 ° C./s or more. An annealing step for holding 1000 s;
The manufacturing method of the high-tensile cold-rolled steel sheet characterized by performing sequentially.   The high-strength cold rolling according to claim 3, wherein the steel slab has a composition that further includes 0.1% or less in total of one or more of Ca and REM in mass%. A method of manufacturing a steel sheet.