JP2010285657A - Precipitation strengthened double-phase cold-rolled steel sheet and method for producing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high strength cold rolled steel sheet having an excellent strength-ductility balance and further having excellent stretch flange formability, and to provide a method for producing the same. <P>SOLUTION: The cold rolled steel sheet has a composition containing, by mass, 0.03 to 0.15% Ti, in which the content of Nb is limited to ≤0.03%, the content of Mo is limited to ≤0.25% and the content of V is limited to ≤0.25%, and satisfying 0.18≤6Ti+25Nb+3Mo+3V≤1.0, and also satisfying 20C+17.1N-5Ti-2.6Nb-2.5Mo-4.7V≥0.6. The Ti-based carbonitrides of 1 to 50 nm are dispersed, the area ratio of ferrite is ≥50%, the area ratio of a hard structure composed of either or both of martensite or bainite or both of them is 5 to 50%, and the total area ratio of the remaining pearlite, residual austenite and cementite is limited to ≤5%. The average grain size of the ferrite is limited to ≤20 μm, and also, the ratio of unrecrystallized ferrite occupied in the ferrite is limited to ≤25%. <P>COPYRIGHT: (C)2011,JPO&INPIT

Description

  The present invention relates to a precipitation-strengthened double-phase cold-rolled steel sheet and a method for producing the same, using strengthening by a hard structure and precipitation strengthening.

  In recent years, in order to improve the fuel consumption of automobiles, the weight of the vehicle body has been reduced. In order to reduce the weight of the vehicle body, it is very effective to reduce the thickness of the member by increasing the strength of the steel plate. On the other hand, many members for automobiles have complicated shapes, and press molding which is low cost is often used for processing these members. Therefore, the steel material is required to have a balance between strength and ductility.

  As a steel material having both high strength and high ductility, a dual phase steel (DP steel: Dual Phase steel) having a composite structure composed of a hard second phase mainly composed of ferrite and martensite is known. However, the DP steel has a large hardness difference between the soft phase ferrite and the hard phase martensite, and microvoids due to the hardness difference of each phase are likely to occur at the boundary. Therefore, local elongation is poor, stretch flangeability is inferior, and it is not suitable for a member subjected to hole expansion processing.

  Therefore, a steel sheet having high strength and high ductility and excellent in stretch flangeability is required. To date, steel sheets in which the grain size of recrystallized ferrite is refined after annealing have been proposed as steel sheets that achieve both high strength and workability (see, for example, Patent Documents 1 to 3).

  However, the cold-rolled steel sheets proposed in Patent Documents 1 and 2 have a very narrow heating temperature range in annealing after cold rolling, and it is extremely difficult to control the temperature of the steel sheet. Furthermore, in the production of the cold-rolled steel sheet proposed in Patent Document 3, it is necessary to immerse the coil after hot rolling in cooling water or to force-cool the coil while rewinding the coil, which impairs productivity. .

  Further, cold rolled steel sheets using precipitation strengthening by dispersion of fine precipitates with emphasis on stretch flangeability have been proposed (see, for example, Patent Documents 4 to 6). The cold-rolled steel sheet proposed in Patent Document 4 has a microstructure of a ferrite single phase and precipitates fine carbides including Ti, Mo, and W. Furthermore, the cold-rolled steel sheets proposed in Patent Documents 5 and 6 precipitate and strengthen ferrite, reduce the hardness difference from the second phase, and improve workability.

Japanese Patent Laid-Open No. 2003-27043 JP 2004-250774 A JP 2005-179732 A JP 2003-321732 A JP 2002-69574 A JP 2007-16319 A

  However, the steel sheet proposed in Patent Document 4 is obtained by refining carbonitride using expensive Mo and W, which increases the cost. Moreover, the steel plates proposed in Patent Documents 5 and 6 are produced by carbonitride during cooling after annealing, and it cannot be said that stretch flangeability is sufficient.

  In view of such circumstances, the present invention provides a cold-rolled steel sheet excellent in strength-ductility balance and stretch flangeability without excessively adding an expensive alloy. It is an object of the present invention to provide a production method for stably and economically obtaining the above without impairing productivity.

The present invention promotes solid solution of Ti before cold rolling, and precipitates fine carbonitride of Ti at the time of continuous annealing after cold rolling, thereby significantly improving the ductility and stretch flangeability of high strength steel sheets. Containing Ti, Nb, Mo, V, which contributes to precipitation strengthening, suppresses the progress of recrystallization, and fixes C and N that contribute to strengthening of the hard structure. The precipitation-strengthening-type duplex steel sheet with controlled amount and the transformation to austenite sufficiently proceed (Ac ₁ +20) ° C., and precipitation strengthened ferrite can be sufficiently secured (Ac ₃ -20) ° C. or less. In order to suppress the formation of pearlite and cementite, the cooling method is performed at a cooling rate of 500 to 650 ° C. at 3 ° C./s or more. The gist of the present invention is as follows.

(1) In mass%,
C: 0.05-0.20%,
Mn: 0.50 to 3.00%,
Ti: 0.03-0.15%,
Containing
Si: 2.50% or less,
Al: 1.50% or less,
P: 0.15% or less,
S: 0.010% or less,
N: 0.0060% or less,
Nb: 0.03% or less (including 0),
Mo: 0.25% or less (including 0),
V: 0.25% or less (including 0)
The content of C, N, Ti, Nb, Mo, and V satisfies the following (formula 1) and the following (formula 2), the balance is composed of Fe and inevitable impurities, The particle size is 1 to 50 nm, the area ratio of ferrite is 50% or more, the area ratio of the hard structure composed of one or both of martensite and bainite is 5 to 50%, the remaining pearlite, residual austenite and The total area ratio of cementite is limited to 5% or less, the average particle diameter of the ferrite is limited to 20 μm or less, and the proportion of the unrecrystallized ferrite in the ferrite is limited to 25% or less. Precipitation strengthened double-phase cold-rolled steel sheet.
0.18 ≦ 6Ti + 25Nb + 3Mo + 3V ≦ 1.0 (Formula 1)
20C + 17.1N-5Ti-2.6Nb-2.5Mo-4.7V ≧ 0.6 (Formula 2)
Here, C, N, Ti, Nb, Mo, and V are contents [mass%] of each element.

(2) By mass%
B: 0.0050% or less,
Cr: 2.00% or less,
Ni: 1.00% or less,
Cu: 1.00% or less,
W: The precipitation strengthening type duplex cold-rolled steel sheet according to (1) above, containing one or more of 0.20% or less.
(3) In mass%,
Ca: 0.010% or less,
REM: The precipitation-strengthened double-phase cold-rolled steel sheet according to (1) or (2) above, containing one or both of 0.100% or less.
(4) The precipitation-strengthened double-phase cold-rolled steel sheet according to any one of (1) to (3), wherein the following (Formula 3) is satisfied.
80- (230C + 20Si + 55Mn-65Al + 125Ti
+ 1760Nb + 480Mo + 210V + 80000B + 35Cr
+ 25Ni + 20Cu + 25W) <0 (Formula 3)
Here, C, Si, Mn, Al, Ti, Nb, Mo, V, B, Cr, Ni, Cu, and W are contents [mass%] of each element.

(5) A precipitation-strengthened double-phase cold-rolled steel sheet, wherein the cold-rolled steel sheet according to any one of (1) to (4) above is hot-dip galvanized.
(6) A precipitation-strengthened double-phase cold-rolled steel sheet, wherein the cold-rolled steel sheet according to any one of (1) to (4) is subjected to alloying hot-dip galvanizing.

(7) A method for producing a cold-rolled steel sheet according to any one of (1) to (4) above, wherein the steel sheet has the chemical component according to any one of (1) to (4) above. Is heated to 1050 ° C. or higher, cooled to 550 ° C. or lower at a cooling rate of 15 ° C./s or higher, and further subjected to cold rolling with a total rolling reduction of 40% or higher, and the maximum heating temperature (Ac ₁ +20) More than + 20 ° C. to (Ac ₃ −20) ° C., a precipitation strengthened double phase cold rolled steel sheet characterized by performing continuous annealing with an average cooling rate from 650 ° C. to 500 ° C. being 3 ° C./s or more. Production method.

(8) A method for producing a cold-rolled steel sheet according to any one of (1) to (4) above, wherein the steel has the chemical component according to any one of (1) to (4) above. The piece is heated to 1050 ° C. or higher, hot-rolled at a finishing temperature of (Ae ₃ −30) ° C. or higher, cooled to 550 ° C. or lower at a cooling rate of 15 ° C./s or more, and the reduction rate And the maximum heating temperature is (Ac ₁ +20) ° C. to (Ac ₃ −20) ° C., and the average cooling rate from 650 ° C. to 500 ° C. is 3 ° C./s. The manufacturing method of the precipitation strengthening type | mold double-phase cold-rolled steel plate characterized by performing the continuous annealing mentioned above.

(9) A method for producing a cold-rolled steel sheet according to (5) above, wherein the cold-rolled steel sheet produced by the production method according to (7) or (8) is subjected to continuous annealing and then into a hot dip galvanizing bath. A method for producing a precipitation-strengthened double-phase cold-rolled steel sheet, characterized by dipping.
(10) A method for producing a cold-rolled steel sheet according to (6) above, wherein the precipitation-strengthening-type composite is characterized in that after being immersed in the hot dip galvanizing bath according to (9), an alloying treatment is performed. A method for producing a phase cold-rolled steel sheet.

  According to the present invention, a precipitation-strengthened cold-rolled steel sheet excellent in strength-ductility balance and stretch flangeability can be obtained economically, without impairing productivity, etc. The contribution is very significant.

It is a figure which shows the relationship between the ratio of the non-recrystallized ferrite which occupies for a ferrite, and intensity-ductility balance. It is a figure which shows the relationship between a precipitation strengthening element and the ratio of the non-recrystallized ferrite which occupies for a ferrite.

  In order to ensure ductility and stretch flangeability and to increase the strength of the steel sheet, it is preferable to generate a hard structure composed of one or both of martensite and bainite and to precipitate and strengthen soft ferrite. Precipitation strengthening is caused by matching strain generated between the hard particles and the surrounding ferrite, and generally the amount of strengthening by the hard particles depends on the density. Therefore, for example, when precipitates such as carbonitrides become coarse and the density decreases, the amount of strengthening becomes small. Therefore, in order to effectively improve the strength of ferrite, a fine Ti system with a particle diameter of 1 to 50 nm It is necessary to disperse the carbonitride.

  However, conventionally, when precipitation strengthening a cold-rolled steel sheet, carbonitride is re-dissolved by annealing and then carbonitride is precipitated again during cooling. Therefore, in the conventional method, carbonitride is precipitated in a state where the dislocation density is reduced by annealing, and it is difficult to refine the carbonitride.

  Therefore, the present inventors precipitated Ti-carbonitrides on recrystallized ferrite (referred to as recrystallized ferrite) during annealing after cold rolling, effectively precipitation strengthening the ferrite, and austenite. A method for securing C and N in a solid phase and forming a multiphase structure was investigated. In the present invention, annealing performed after cold rolling is important for precipitating fine Ti-based carbonitrides on recrystallized ferrite and securing a hard structure composed of one or both of martensite and bainite. .

In order to ensure stable austenite that _first undergoes ferrite recrystallization and Ti carbonitride precipitation at the same time and secures a stable austenite that becomes a hard structure after cooling, the present inventors firstly exceed (Ac ₁ +20) ° C. It has been found that it is necessary to heat to (Ac ₃ -20) ° C. This is because Ti carbonitride precipitates preferentially on ferrite and Ti carbonitride does not precipitate on austenite for the following reasons.

  In the two-phase region where austenite and ferrite coexist, ferrite recrystallization is suppressed. When the recrystallization of ferrite is suppressed, the dislocation density increases, and Ti-based carbonitrides are finely precipitated using the dislocations as precipitation nuclei. On the other hand, austenite has a large solid solution amount of C and N, and precipitation of Ti-based carbonitride is suppressed as compared with ferrite. Therefore, after cooling, it becomes a multiphase structure of ferrite in which Ti-based carbonitride is finely precipitated and a hard structure composed of one or both of martensite and bainite.

However, when the annealing temperature exceeds (Ac ₃ -20) ° C., austenite increases and the concentrations of C and N decrease. For this reason, austenite becomes unstable, and after cooling, it transforms into ferrite and cannot secure a hard structure. Ferrite transformed from austenite (called transformed ferrite) is soft because Ti carbonitride does not precipitate. As a result, as with the conventional multiphase steel sheet, the ductility is improved, but the stretch flangeability is lowered.

  Furthermore, if a precipitation strengthening element is added excessively in order to improve the strength, recrystallization of ferrite becomes insufficient during annealing after cold rolling. Moreover, since C and N are fixed, austenite becomes unstable, and after cooling, it transforms into ferrite and cannot secure a hard structure. In particular, when non-recrystallized ferrite remains, the strength-ductility balance decreases. This is because the dislocation density of the non-recrystallized ferrite in which the dislocation substructure called subgrains formed by recovery is high and hard, and the ductility deteriorates due to the precipitation of fine Ti-based carbonitrides. It is. Non-recrystallized ferrite has a structure in which the strain introduced by cold rolling is not recrystallized by annealing and recovered.

  Therefore, the inventors examined the relationship between the strength-ductility balance of the cold-rolled steel sheet and the area ratio of non-recrystallized ferrite in the ferrite. Tensile test pieces according to JIS Z 2201 were collected from various cold-rolled steel sheets, and the tensile test was performed according to JIS Z 2241. Further, the microstructure of the steel sheet was observed with an optical microscope. The area ratio was 50% or more of ferrite, the hard structure was 5 to 50%, and the total area ratio of the remaining pearlite, residual austenite and cementite was 5%. It was confirmed that:

  Among the ferrites, the proportion of unrecrystallized ferrite was measured by analyzing an electron beam backscattering diffraction pattern (EBSP). Discrimination between unrecrystallized ferrite and other ferrites, that is, recrystallized ferrite and transformed ferrite, was performed by analyzing crystal orientation measurement data of EBSP by the Kernel Average Misoration method (KAM method).

  The results are shown in FIG. The vertical axis in FIG. 1 is the product of tensile strength (TS) and elongation at break (El), and the horizontal axis is the proportion of unrecrystallized ferrite in the ferrite. As can be seen from FIG. 1, when the proportion of non-recrystallized ferrite exceeds 25%, the strength-ductility balance is significantly deteriorated. Therefore, in order to improve the strength-ductility balance, it is necessary to limit the proportion of unrecrystallized ferrite in the ferrite to 25% or less.

  In order to suppress the formation of unrecrystallized ferrite, it is necessary to limit the addition of an element that delays recrystallization. However, the precipitation strengthening element contributing to the strengthening of ferrite is also an element that delays recrystallization. Therefore, the present inventors have examined the relationship between the amount of Ti, Nb, Mo, and V that contribute to precipitation strengthening and the formation of unrecrystallized ferrite.

  The results are shown in FIG. The horizontal axis in FIG. 2 is an empirical formula in which the content of each element is weighted according to the strength of the influence of Ti, Nb, Mo, and V on recrystallization. FIG. 2 shows that when 6Ti + 25Nb + 3Mo + 3V is 1.0 or less, the proportion of unrecrystallized ferrite is 25% or less.

  Furthermore, the present inventors also examined the contribution to Ti, Nb, Mo, and V element precipitation strengthening. As a result, the addition of Ti is the most effective for ferrite precipitation strengthening, and in order to form the fine Ti carbonitride necessary for precipitation strengthening, 0.03% or more of Ti is necessary. It was found that the content of Nb, Mo, V may be 0 when the Ti content is 0.03% or more.

Therefore, in the present invention, Ti may be contained alone, or one or more of Nb, Mo, and V may be contained in addition to Ti. When these elements are contained together with Ti, the content [% by mass] must satisfy the following (formula 1). In addition, when Nb, Mo, and V are not added intentionally, it is calculated as 0 in (Formula 1).
0.18 ≦ 6Ti + 25Nb + 3Mo + 3V ≦ 1.0 (Formula 1)

  Moreover, in this invention, in order to fully raise the intensity | strength of a steel plate, it is necessary to ensure the hard structure which consists of one or both of a martensite and a bainite, and to raise the intensity | strength of a hard structure. The hard structure increases as the amount of solid solution C and N in the austenite generated during annealing increases, and the strength increases. On the other hand, precipitation strengthening elements Ti, Nb, Mo, and V generate carbonitrides during annealing. Therefore, in order to secure C and N dissolved in austenite, consider the consumption of C by carbonitrides. There is a need to.

From this viewpoint, the present inventors have studied and found that the steel composition must satisfy the following (formula 2) in order to ensure the strength of the hard structure.
20C + 17.1N-5Ti-2.6Nb-2.5Mo-4.7V ≧ 0.6
... (Formula 2)
Here, C, N, Ti, Nb, Mo, and V are contents [mass%] of each element.

  The above (formula 2) is an empirical formula in which the content of each element is weighted according to the contribution to the stabilization of C, N austenite and the ability to produce carbonitrides of Ti, Nb, Mo, V. It is. A steel sheet satisfying the above (Formula 2) has a sufficient amount of solid solution C in the austenite after the two-phase region annealing, and can secure a hard structure having sufficient strength even after cooling.

  Further, in order to precipitate fine Ti-based carbonitrides in ferrite, it is necessary to deposit carbonitrides during heating after annealing after cold rolling. For this purpose, the precipitation strengthening element must be dissolved in the cold-rolled steel sheet before heating, that is, the hot-rolled steel sheet. For example, it is preferable that the hot-rolled steel sheet is heated to 1050 ° C. or more to dissolve the carbonitride, and then water-cooled to 550 ° C. or less at which the precipitation rate decreases.

  However, in order to reduce the manufacturing cost, it is necessary to increase the cooling rate at a temperature at which Ti-based carbonitride is likely to precipitate when the steel slab is hot-rolled into a steel plate. Therefore, in the present invention, after hot rolling, it is cooled to 550 ° C. or less, wound up, and a precipitation strengthening element such as Ti is dissolved in the hot rolled steel sheet.

Further, during cooling after hot rolling, carbonitrides are easily generated in ferrite transformed from austenite. Therefore, it is preferable to delay the progress of the ferrite transformation. That is, by adding an element that improves hardenability and lowering the Ar ₃ transformation point, precipitation of Ti-based carbonitrides on the hot-rolled steel sheet can be suppressed.

In order to lower the Ar ₃ transformation point and to secure the solid solution Ti amount in the hot-rolled steel sheet, it is preferable to satisfy the following (formula 3). The following (Equation 3) is a result of a study conducted by the inventors based on an equation obtained by quantifying the influence of each element on the Ar ₃ transformation point.

80- (230C + 20Si + 55Mn-65Al + 125Ti
+ 1760Nb + 480Mo + 210V + 80000B + 35Cr
+ 25Ni + 20Cu + 25W) <0 (Formula 3)
Here, C, Si, Mn, Al, Ti, Nb, Mo, V, B, Cr, Ni, Cu, and W are contents [mass%] of each element.

  Hereinafter, the present invention will be described in detail.

  First, the components of the steel sheet of the present invention will be described. In the following description of the components,% means mass%.

C: 0.05-0.20%
C is an element that contributes to an increase in strength. In the present invention, C is an extremely important element for precipitation strengthening with a fine Ti-based carbonitride and securing and strengthening of a hard structure. If the amount of C is less than 0.05%, the area ratio of the hard structure decreases, and sufficient precipitation strengthening cannot be obtained. In order to ensure the strength of the hard tissue, the C content is preferably 0.07% or more. On the other hand, if C is added in excess of 0.20%, coarse iron-based carbides are produced in the steel sheet after annealing and the ductility is impaired, so the upper limit is made 0.20%. In this respect, the amount to be added is preferably 0.15% or less, and more preferably 0.12% or less.

Mn: 0.50 to 3.00%
Mn is an element that lowers the Ar ₃ transformation point, which is the temperature at which austenite transforms into ferrite when heated to a high temperature during hot rolling. When the transformation to ferrite during hot rolling is delayed by the addition of Mn, it becomes easy to secure solid solution Ti in the hot-rolled steel sheet. Mn is an important element for improving the hardenability and obtaining a sufficient amount of hard structure. In order to obtain this effect, it is necessary to add 0.50% or more of Mn. In order to lower the Ar ₃ transformation point, the Mn content is preferably 0.80% or more, and more preferably 1.00% or more.

  On the other hand, if the amount of Mn added exceeds 3.00%, ductility decreases due to segregation, so the upper limit is made 3.00%. In addition, when hot dip galvanizing is performed, if the amount of Mn exceeds 2.50%, the adhesion of galvanizing may be hindered, so the addition amount is preferably 2.50% or less.

Ti: 0.03-0.15%
Ti is the most important element in the present invention. If the amount of Ti added is less than 0.03%, the amount of Ti-based fine carbonitrides with an average particle size of 1 to 50 nm that contributes to precipitation strengthening is small, and a sufficient amount of strengthening cannot be obtained. The lower limit. From this viewpoint, it is preferable that the addition amount be 0.04% or more. On the other hand, when Ti is added excessively, recrystallization during annealing is suppressed, and ductility deteriorates. On the other hand, if the added amount exceeds 0.15%, the effect of increasing the strength is saturated. From the above viewpoint, the upper limit of the addition amount is 0.15%. In order to ensure ductility, the amount of Ti added is preferably 0.12% or less, and more preferably 0.10% or less.

Si: 2.50% or less Si is a deoxidizing element, and the lower limit is not particularly defined. However, if the addition amount is less than 0.001%, the production cost is increased, so 0.001% or more is preferable. Further, since Si is an element that increases the strength by solid solution strengthening, addition of 0.1% or more is preferable. Furthermore, since Si raises the Ac ₁ point, annealing at a high temperature is possible, and it becomes easy to ensure ductility by promoting recrystallization. From these viewpoints, Si may be positively added according to the target strength level. However, if the added amount exceeds 2.50%, the ductility deteriorates greatly, so the Si amount is 2.50%. Restrict to:

  Moreover, since chemical conversion processability will fall when there is much Si amount, Si amount is preferably 1.50% or less. Furthermore, when hot dip galvanizing is performed, problems such as a decrease in plating adhesion and a decrease in productivity due to a delay in the alloying reaction may occur, so the Si content is preferably 1.20% or less.

Al: 1.50% or less Al is a deoxidizing element, and the lower limit is not particularly limited, but is preferably 0.01% or more from the viewpoint of deoxidation. Further, Al is an element that remarkably increases the Ac ₁ point, enables annealing at a high temperature, and facilitates ensuring ductility by promoting recrystallization. However, excessive Al addition significantly increases the Ar ₃ point, and if the addition amount exceeds 1.50%, hot rolling at a temperature (γ region) where the structure is austenite becomes difficult, so the Al amount is 1 Limit to 50% or less.
In addition, when hot dip galvanizing is performed, problems such as a decrease in plating adhesion and a decrease in productivity due to a delay in the alloying reaction may occur, so that the Al addition amount is preferably 1.20% or less.

P: 0.15% or less P is an impurity. However, since P shows a large solid solution strengthening, P may be positively added when the strength needs to be increased. However, if the addition amount exceeds 0.15%, fatigue strength after spot welding and secondary workability deteriorate. Therefore, the P content is limited to 0.15% or less. Furthermore, since the alloying reaction is extremely slow during hot dip galvanizing and the productivity is reduced, the P content is preferably 0.05% or less.

S: 0.010% or less S is an impurity, and if the content exceeds 0.010%, it causes hot cracking and deterioration of workability. Therefore, the amount of S is limited to 0.010% or less.

N: 0.0060% or less N is an element inevitably contained in the steel. When the content exceeds 0.0060%, TiN is generated and ductility is impaired. Therefore, the N content is limited to 0.0060% or less. Moreover, when coarse TiN is produced, Ti contributing to precipitation strengthening decreases, so the N content is preferably 0.0050% or less, and more preferably 0.0040% or less. The lower limit of the N amount is not particularly set, but setting the content to less than 0.0005% brings about an increase in manufacturing cost, so the lower limit is preferably set to 0.0005%.

Nb: 0.03% or less Nb is an element that forms fine carbonitrides with Ti and contributes to precipitation strengthening, and is selectively added. However, Nb remarkably delays recrystallization during annealing, so adding excessive Nb degrades ductility. Therefore, the Nb content is limited to 0.03% or less.

Mo: 0.25% or less Mo is mainly carbide is formed, is an element contributing to precipitation strengthening are selectively added. However, Mo, like Nb, remarkably delays recrystallization during annealing, so the ductility deteriorates when excessive Mo is added. Therefore, the Mo content is limited to 0.25% or less.

V: 0.25% or less V is an element that forms fine carbonitrides like Nb and contributes to precipitation strengthening, and is selectively added. However, V also significantly delays recrystallization during annealing, so the content is limited to 0.25% or less in order to suppress deterioration of ductility.

(Formula 1)
Furthermore, in order to obtain a good strength-ductility balance, it is necessary to satisfy the following (formula 1) particularly for Ti, Nb, Mo, and V. Thereby, delay of recrystallization can be prevented and fine Ti-based carbonitride can be precipitated during annealing. In the following (Formula 1), Ti, Nb, Mo, and V are the contents [% by mass] of each element. When Nb, Mo, and V are not intentionally contained, 0 is calculated.
0.18 ≦ 6Ti + 25Nb + 3Mo + 3V ≦ 1.0 (Formula 1)

(Formula 2)
Furthermore, in order to ensure the amount and strength of the hard structure of the cold-rolled steel sheet after annealing, the contents of C, N, Ti, Nb, Mo, and V need to satisfy the following (formula 2). . When the following (Formula 2) is satisfied, when annealing in the two-phase region of austenite and ferrite, C and N that are dissolved in austenite are secured, and a hard structure having sufficient strength after cooling is secured. Can do. In the following (Formula 2), Ti, Nb, Mo, and V are the content [% by mass] of each element. When Nb, Mo, and V are not intentionally contained, 0 is calculated.
20C + 17.1N-5Ti-2.6Nb-2.5Mo-4.7V ≧ 0.6
... (Formula 2)

Furthermore, in order to promote solid solution of Ti to the hot-rolled steel sheet, one or more of B, Cr, Ni, Cu, and W that lower the Ar ₃ transformation point may be added.

B: 0.0050% or less B is an element that significantly lowers the Ar ₃ transformation point. In order to delay the transformation to ferrite during cooling after hot rolling and to secure solid solution Ti of the hot-rolled steel sheet, 0.0002% or more is preferably added. In order to enhance this effect, it is more preferable to add B in an amount of 0.0005% or more. On the other hand, when B is added excessively, the effect is saturated and the workability is deteriorated, so the B content is preferably made 0.0050% or less.

Cr: 2.00% or less Cr is an element that lowers the Ar ₃ transformation point. In order to secure solid solution Ti of the hot-rolled steel sheet, it is preferable to add 0.01% or more of Cr. Cr is an element that increases the Ac ₁ point and enables annealing at a high temperature. In order to ensure ductility by promoting recrystallization, it is preferable to add 0.05% or more of Cr. On the other hand, when Cr is added excessively, workability is impaired, so the Cr content is preferably made 2.00% or less.

Ni: 1.00% or less Ni is an element that lowers the Ar ₃ transformation point. To secure solid solution Ti of the hot-rolled steel sheet, it is preferable to add 0.01% or more of Ni. Further, Ni is more preferably added in an amount of 0.05% or more in order to enhance this effect. On the other hand, Ni is an expensive element, and if it is added excessively, the manufacturing cost increases, so the upper limit is preferably made 1.00% or less.

Cu: 1.00% or less Cu is an element that lowers the Ar ₃ transformation point, and in order to secure solid solution Ti of the hot-rolled steel sheet, addition of 0.01% or more of Cu is preferable. Further, Cu is an element that contributes to strengthening of steel, and addition of 0.05% or more is preferable. On the other hand, it is preferable to make Cu amount into 1.00% or less from a viewpoint of economical efficiency.

W: 0.20% or less W is an element that lowers the Ar ₃ transformation point, and in order to secure solid solution Ti of the hot-rolled steel sheet, 0.01% or more of W is preferably added. W is an element that forms fine carbonitride with Ti and contributes to precipitation strengthening, and is preferably added in an amount of 0.05% or more. On the other hand, from the viewpoint of economy, the W amount is set to 0.20% or less.

(Formula 3)
When the ferrite transformation proceeds at a high temperature after hot rolling, Ti carbonitrides are generated inside the ferrite, and the solid solution Ti in the steel sheet is greatly reduced. In order to avoid this, it is preferable to satisfy the following (formula 3).
80- (230C + 20Si + 55Mn-65Al + 125Ti
+ 1760Nb + 480Mo + 210V + 80000B + 35Cr
+ 25Ni + 20Cu + 25W) <0 (Formula 3)
Here, C, Si, Mn, Al, Ti, Nb, Mo, V, B, Cr, Ni, Cu, and W are contents [mass%] of each element. The steel sheet that satisfies the above (Formula 3) has sufficient hardenability after hot rolling, and can easily secure solid solution Ti in the steel sheet. In particular, when the average cooling rate from the completion of hot rolling until the steel sheet temperature reaches 550 ° C. is smaller than 20 ° C./s, it is preferable to satisfy the above (Formula 3).

  In order to control the form of inclusions and improve the material, one or both of Ca and REM may be contained as necessary.

Ca: 0.010% or less Ca is an element that mainly contributes to improvement of the material by generating sulfides and controlling the form of inclusions. In order to obtain the effect, it is preferable to add 0.0005% or more of Ca. However, if added excessively, the ductility may be impaired, so the Ca content is preferably 0.010% or less.

REM: 0.10% or less REM is an element that generates oxides and sulfides, controls the form of inclusions, and contributes to the improvement of the material. In order to obtain the effect, it is preferable to add REM 0.0005% or more. However, if added excessively, ductility may be impaired, so the REM content is preferably made 0.10% or less.

  Inevitable impurities include, for example, Sn, and the effect of the present invention is not impaired even if an element mixed from scrap is contained in the range of 0.1% by mass or less.

  Next, the microstructure of the steel sheet of the present invention will be described.

  The microstructure in the steel sheet of the present invention is composed of ferrite, a hard structure composed of one or both of martensite and bainite, and one or more of remaining pearlite, retained austenite, and cementite.

  The ferrite area ratio is 50% or more in order to ensure the ductility of the steel sheet. The area ratio of the hard structure composed of one or both of martensite and bainite is 5% or more in order to increase the strength of the steel sheet. On the other hand, if the hard structure increases too much, the ductility deteriorates, so the area ratio of the hard structure is set to 50% or less. The balance consists of one or more of pearlite, retained austenite, and cementite other than ferrite and hard structure, and these deteriorate the stretch flangeability, so the total area ratio is limited to 5% or less.

  The microstructure is obtained by taking a sample with the cross section of the plate thickness parallel to the rolling direction as the observation surface, polishing the observation surface, performing nital etching, if necessary, repeller etching or picral etching, and observing with an optical microscope. The area ratio of the ferrite and the hard structure can be obtained by analyzing the image of the microstructure photograph obtained by the optical microscope. Note that the area ratio obtained from the microstructure is the same as the volume ratio.

  The ferrite grain size greatly affects the properties of the steel sheet. If the ferrite grain size is coarse, the strength of the steel sheet cannot be secured, so 20 μm is the upper limit. From this viewpoint, the thickness is preferably 15 μm or less, and more preferably 10 μm or less. The lower limit of the ferrite grain size is not particularly set, but if it is less than 1.0 μm, the cost is increased and the ductility is greatly impaired. From the viewpoint of productivity, the ferrite particle size is more preferably 1.5 μm or more.

The particle diameter of the ferrite can be obtained, for example, by observing an arbitrary portion of the steel sheet using an optical microscope, measuring the number of ferrite grains in the range of 1000 μm ² or more, and determining the average equivalent circle diameter.

  Ferrite is a general term for recrystallized ferrite, transformed ferrite, and non-recrystallized ferrite. In addition, in the structure observation with an optical microscope, the difference between recrystallized ferrite and transformed ferrite is not clear, and it is difficult to distinguish them. If the proportion of unrecrystallized ferrite in the ferrite increases, the strength increases, but the ductility deteriorates greatly, so this is limited to 25% or less. From this viewpoint, it is preferably 20% or less, more preferably 10% or less.

  Unrecrystallized ferrite and other ferrites, that is, recrystallized ferrite and transformed ferrite can be discriminated by analyzing the crystal orientation measurement data of EBSP by the KAM method.

  In the grains of unrecrystallized ferrite, although dislocations are recovered, there is a continuous change in crystal orientation caused by plastic deformation during cold rolling. On the other hand, in the recrystallized ferrite and the transformed ferrite, the crystal orientation change in the ferrite grains is extremely small. This is because although the crystal orientation of adjacent crystal grains varies greatly due to recrystallization and transformation, the crystal orientation hardly changes within one crystal grain.

  The KAM method can quantitatively indicate the crystal orientation difference between adjacent pixels (measurement points). Accordingly, in the present invention, when the average crystal orientation difference between adjacent measurement points is within 1 ° and the pixel boundary is defined as a grain boundary between the pixels having the average crystal orientation difference of 5 ° or more, the crystal grain size exceeds 0.5 μm. Grains are defined as ferrite other than unrecrystallized ferrite, that is, recrystallized ferrite and transformed ferrite.

The EBSP measurement may be performed within a range of 10000 μm ² or more at a measurement position of a thickness of ¼ of any plate cross section at a measurement interval of 0.5 μm or less. As a result of the EBSP measurement, the measurement points obtained are output as pixels. The sample to be used for EBSP crystal orientation measurement is a method in which the steel plate is reduced to a predetermined thickness by mechanical polishing or the like, and then the strain near the observation surface is removed by electrolytic polishing or the like. Create as follows.

  The sample used for measuring the crystal orientation of EBSP is nital etched, the optical microscope photograph of the same field of view as the measurement of EBSP is measured at the same magnification, and the obtained structure photograph is image-analyzed. The proportion of ferrite can be determined.

The fine Ti-based carbonitride contained in the ferrite improves the strength of the ferrite without greatly impairing the ductility. If the average particle size of the Ti-based carbonitride is large, sufficient strengthening cannot be obtained, so the upper limit is made 50 nm. The lower limit of the average particle size of the Ti-based carbonitride is not particularly specified, but the strengthening amount is saturated at 1 nm or less. This carbide is mainly composed of Ti and C, but may contain elements such as Nb, Mo, V, Cr, Fe, and N. Specifically, Ti (C, N), Ti (C, N), Nb (C, N), V (C, N), Fe ₃ C, Cr ₂₃ C ₆ , Cr ₂ N, Examples include composite precipitates with Mo ₂ C.

  The method for measuring the particle size of the Ti-based carbonitride is not particularly specified. For example, an extraction replica sample is prepared from an arbitrary portion of the steel plate, and observed at a magnification of 10,000 times or more using a transmission electron microscope (TEM). It is obtained by calculating the average equivalent circle diameter of the carbonitride. Ti-based carbonitrides are particles in which Ti is detected by an energy dispersive spectrometer (EDS) attached to the TEM.

  In this invention, when there is little solid solution Ti amount of the cold rolled steel plate before annealing, sufficient precipitation strengthening is not obtained after annealing, but intensity | strength falls. Therefore, the amount of solute Ti in the hot-rolled steel sheet before cold rolling is preferably 0.01% or more, and more preferably 0.02% or more. In addition, the amount of solid solution Ti in a steel plate is calculated | required by measuring the amount of non-solution Ti in a steel plate by methods, such as an extraction residue method, and subtracting from the amount of addition Ti.

  Next, the manufacturing method of the steel plate of this invention is demonstrated.

  The cold-rolled steel sheet of the present invention is manufactured by cold-rolling and annealing a hot-rolled steel sheet in which solute Ti is secured. The raw material before cold rolling is manufactured by heat-treating a hot-rolled steel sheet manufactured by a conventional method, or by hot rolling in which a finishing temperature, a cooling rate, and a winding temperature are controlled.

  In order to heat-treat a hot-rolled steel sheet produced by a conventional method and obtain a steel sheet having a sufficient amount of dissolved Ti, it is necessary to set the heating temperature of the heat treatment to 1050 ° C. or higher. This is for sufficiently dissolving the Ti carbonitride deposited on the hot-rolled steel sheet, and the preferred heating temperature is 1100 ° C. or higher.

  After heating to the heat treatment temperature, accelerated cooling such as water cooling is performed. In order to suppress the precipitation of Ti-based carbonitrides, it is necessary to set the cooling rate of accelerated cooling to 15 ° C./s or more. In addition, a cooling rate is an average cooling rate until the temperature of a steel plate will be 550 degreeC after acceleration cooling is started after completion of finish rolling. The cooling rate is preferably 20 ° C./s or more, and more preferably 25 ° C./s or more.

  The stop temperature for accelerated cooling is set to 550 ° C. or lower in order to suppress precipitation of Ti-based carbonitride. In order to prevent the precipitation of Ti-based carbonitrides, it is preferable to stop accelerated cooling at 450 ° C. or less, and it is more preferable to set the stop temperature to 350 ° C. or less.

  From the viewpoint of productivity, it is preferable to manufacture a steel plate having a sufficient amount of solute Ti in the hot rolling process, not by the above heat treatment. In this case, the steel slab is heated and hot-rolled, and then it is wound with accelerated cooling in order to suppress the precipitation of Ti-based carbonitride.

  The steel piece to be subjected to hot rolling may be melted and cast by a conventional method. Although this steel slab may be a forged or rolled steel ingot, it is preferable to manufacture the steel slab by continuous casting from the viewpoint of productivity. Moreover, you may manufacture with a thin slab caster. Typically, the steel slab is cooled after casting and reheated for hot rolling.

  The heating temperature at the time of hot rolling is set to 1100 ° C. or higher in order to dissolve the coarse Ti-based carbonitride generated during casting and secure sufficient solute Ti. In order to heat the steel piece efficiently and uniformly, the heating temperature is preferably 1150 ° C. or higher. The upper limit of the heating temperature is not particularly defined, but when heated above 1300 ° C., the crystal grain size of the steel slab becomes coarse and the workability may be impaired. Also, a process such as continuous casting-direct rolling (CC-DR) in which hot rolling is performed immediately after casting the molten steel may be employed.

The finishing temperature of hot rolling is (Ae ₃ -30) ° C. or higher. Ae ₃ is the transformation temperature in the equilibrium state. When the finishing temperature of hot rolling is set to a temperature lower than (Ae ₃ -30) ° C., ferrite is generated during rolling or immediately after rolling, precipitation of Ti carbonitride is promoted, and sufficient solid solution Ti is obtained. This is because it cannot be obtained.

  After completion of hot rolling, accelerated cooling is performed. When the cooling rate of accelerated cooling is low, precipitation of Ti-based carbonitride proceeds during cooling, and sufficient solute Ti cannot be obtained. Therefore, a cooling rate shall be 15 degrees C / s or more. In order to ensure solid solution Ti, the cooling rate is preferably 20 ° C./s or more, and more preferably 25 ° C./s or more.

  After accelerated cooling, the hot rolled steel sheet is wound up. When the coiling temperature after hot rolling is high, precipitation of Ti carbonitride proceeds during cooling, and sufficient solid solution Ti cannot be obtained. Therefore, the upper limit of the coiling temperature is set to 550 ° C. In this respect, the winding temperature is preferably 450 ° C. or lower, and more preferably 350 ° C. or lower.

  The hot-rolled steel sheet heat-treated as described above or the hot-rolled steel sheet that has been hot-rolled and cooled under controlled conditions is cold-rolled and annealed.

  When the rolling reduction of cold rolling is small, recrystallization of ferrite does not proceed sufficiently during annealing, and ductility is greatly deteriorated. Therefore, the lower limit of the cold rolling rate is 40%. In order to refine the crystal grain size of the cold rolled steel sheet, the cold rolling rate is preferably 50% or more, and more preferably 60% or more. On the other hand, the upper limit of the cold rolling rate is not particularly defined, but it is preferably 90% or less from the viewpoint of productivity.

  In the present invention, annealing after cold rolling is important. Annealing is preferably performed by continuous annealing equipment in order to control the heating temperature.

If the annealing temperature is too low with respect to the Ac ₁ transformation point, austenite is not generated and a single-phase structure of recrystallized ferrite is obtained. Therefore, the lower limit of the heating temperature is over (Ac ₁ +20) ° C. On the other hand, when the heating temperature exceeds (Ac ₃ -20) ° C., austenite increases, the amount of solid solution C and N decreases, and austenite becomes unstable. As a result, the transformed ferrite is generated after cooling, the strength cannot be secured, and the stretch flangeability is also deteriorated. Therefore, the upper limit of the heating temperature is (Ac ₃ -20) ° C. or less.

  The heating rate during annealing is not particularly specified, but in order to proceed with recrystallization, it is preferable to reduce the heating rate of annealing to 100 ° C./s or less. In order to promote recrystallization and reduce the area ratio of unrecrystallized ferrite, the heating rate is preferably 50 ° C./s or less, more preferably 20 ° C./s or less. On the other hand, when the heating rate is less than 0.1 ° C./s, Ti-based carbonitride grows and the strength may be lowered. Therefore, the heating rate is preferably set to 0.1 ° C./s or more.

  Moreover, when the cooling rate after annealing is slow, austenite is transformed into pearlite and cementite, and a hard structure cannot be obtained. Therefore, the average cooling rate between 650 ° C. and 500 ° C. at which pearlite and cementite are easily generated is set to 3 ° C./s or more. In order to suppress the formation of pearlite and cementite, the cooling rate between 500 and 650 ° C. is preferably 5 ° C./s or more, and more preferably 10 ° C./s or more. On the other hand, in order to make the cooling rate over 500 ° C./s, it is necessary to introduce special equipment and the like, so it is preferable to set the upper limit of the cooling rate to 500 ° C./s. The cooling rate after annealing may be appropriately controlled by forced cooling with water or the like, blowing of refrigerant, blowing air, mist, or the like.

  After annealing, hot dip galvanization or galvannealing may be performed as necessary. The composition of the galvanizing is not particularly limited, and besides zinc, Fe, Al, Mn, Cr, Mg, Pb, Sn, Ni, etc. may be added as necessary.

  The alloying treatment is preferably performed in a temperature range of 450 to 600 ° C. after hot dip galvanization. This is because alloying does not proceed sufficiently at temperatures below 450 ° C., and alloying proceeds excessively at temperatures above 600 ° C., and the plating layer becomes brittle. This is to induce problems. The alloying time is 5 s or longer. If it is less than 5 s, alloying does not proceed sufficiently. The upper limit is not particularly defined, but is preferably 100 s or less from the viewpoint of production efficiency.

  After annealing, pickling may be performed as necessary. Further, the cold-rolled steel sheet, hot-dip galvanized steel sheet, and alloyed hot-dip galvanized steel sheet may be subjected to skin pass rolling inline or offline. In order to force the shape, the rolling reduction of the skin pass rolling is preferably 0.1% or more. On the other hand, if the rolling reduction of the skin pass exceeds 5.0%, the workability may be impaired.

  Furthermore, the cold-rolled steel sheet and the various plated steel sheets can be subjected to a surface treatment such as coating with an organic film, an inorganic film, or coating with various paints depending on the purpose.

  Steel pieces having the composition shown in Table 1 were melted to produce steel pieces. A hot-rolled steel sheet produced from this steel slab by a conventional method was re-heated under the conditions shown in Table 2 and then water-cooled to 550 ° C. or lower, cold-rolled and continuously annealed to obtain a cold-rolled steel sheet.

Formula 1 of Table 1 is the value of the left side of the following (Formula 1) calculated by the content [mass%] of each element of Ti, Nb, Mo, and V.
6Ti + 25Nb + 3Mo + 3V ≦ 1.0 (Formula 1)

Formula 2 in Table 1 is a value on the left side of the following (Formula 2) calculated by the content [% by mass] of each element of C, N, Ti, Nb, Mo, and V.
20C + 17.1N-5Ti-2.6Nb-2.5Mo-4.7V ≧ 0.6
... (Formula 2)

Moreover, Formula 3 of Table 1 is the following (formula calculated by content [mass%] of each element of C, Si, Mn, Al, Ti, Nb, Mo, V, B, Cr, Ni, Cu, and W. The value on the left side of 2).
80- (230C + 20Si + 55Mn-65Al + 125Ti
+ 1760Nb + 480Mo + 210V + 80000B + 35Cr
+ 25Ni + 20Cu + 25W) <0 (Formula 3)
In addition, Formulas 1-3 are the same also in Table 4.

The heating temperature [° C.] in Table 2 is the reheating temperature of the steel sheet. The cold rolling rate is a value obtained by dividing the difference between the plate thickness before the start of cold rolling and the plate thickness after completion by the plate thickness before the start of cold rolling, and is expressed as a percentage. HF [° C.] is the maximum heating temperature in the continuous annealing process. Ac ₁ [° C.] and Ac ₃ [° C.] were determined by measuring the volume change of the test piece during heating at 10 [° C./s] using a thermal dilatometer. CRc [° C./s] is an average cooling rate between 650 ° C. and 500 ° C. in cooling after annealing.

  The microstructure of these steel sheets was observed with an optical microscope, image analysis was performed, the area ratio of ferrite bainite and martensite was determined, and the crystal grain size of ferrite was measured. The ratio of non-recrystallized ferrite in the ferrite was determined from an optical microscopic structure photograph by EBSP measurement and analysis by the KAM method. Furthermore, a replica sample was prepared, TEM observation was performed, and the particle size of the Ti carbonitride was measured.

  Moreover, the tensile test piece based on JISZ2201 was extract | collected from the obtained steel plate, the tensile test was done based on JISZ2241, and the tensile strength was measured. The hole expansion test was evaluated according to the test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996.

Table 3 shows the microstructure observation results and materials of the steel sheet. In Table 3, Va [%] is the area ratio of ferrite, Vb [%] and Vm [%] are the area ratios occupied by bainite and martensite, respectively, and the hard structure [%] is one or both of martensite and bainite. Is an area ratio occupied by a hard structure composed of Vb and Vm. V _other is an area ratio occupied by a tissue other than the above, and was obtained by the following (formula 4).
V _other [%] = 100−Va−Vb−Vm (Formula 4)

  Non-recrystallized α [%] is the ratio of non-recrystallized ferrite in the ferrite, and d (ferrite) [μm] is the ferrite particle diameter, indicating the equivalent circle diameter determined by image analysis. d (precipitate) [nm] is the particle diameter of the Ti-based carbonitride, and indicates the equivalent-circle diameter of the particles obtained by analyzing the TEM photograph. Also, λ [%] is the hole expansion rate.

  As is apparent from Table 3, when the steel having the chemical component of the present invention is cold-rolled and annealed under appropriate conditions, it can be seen that both excellent strength-ductility balance and excellent stretch flangeability are achieved. .

  In addition to the steel slabs listed in Table 1, steel having the composition shown in Table 4 was melted to produce steel slabs. The blank in Table 4 means that no element was intentionally added. Moreover, the steel strips shown in Table 1 and Table 4 were subjected to hot rolling, cold rolling and continuous annealing under the conditions shown in Tables 5 and 6 to obtain cold-rolled steel sheets. Example 1 shows the area ratio of ferrite, bainite and martensite of the steel sheet, the crystal grain size of ferrite, the ratio of unrecrystallized ferrite in the ferrite, the particle diameter of Ti-based carbonitride, the tensile test and the hole expansion test. And performed in the same manner.

SRT [° C.] in Tables 5 and 6 is the heating temperature of the steel slab before hot rolling. FT [° C.] is the finishing temperature of hot rolling, and is the temperature of the steel sheet measured after the final pass of rolling, that is, on the finishing side. CR [° C./s] is the average cooling rate from the completion of finish rolling until the steel plate temperature reaches 550 ° C., and CT is the coiling temperature. Ae ₃ [° C.] was determined from the composition of each steel by thermodynamic calculation. Tables 7 and 8 show the microstructure observation results and materials of the steel plates evaluated in the same manner as in Example 1. In Tables 4 to 8, the underline means outside the scope of the present invention or outside the preferred range.

  As is clear from Tables 7 and 8, when the steel having the chemical composition of the present invention is hot-rolled, cold-rolled and annealed under appropriate conditions, both excellent strength-ductility balance and excellent stretch flangeability are achieved. You can see that

  On the other hand, production No. Nos. 55-61 are steel Nos. Whose steel compositions are outside the scope of the present invention. It is a comparative example using TZ. Production No. No. 55 is an example in which the strength is lowered because the amount of C is small and the precipitation strengthening of ferrite and the strengthening by the hard structure become insufficient. Production No. 56 is an example in which the amount of Ti is small, the precipitation strengthening of ferrite becomes insufficient, and the stretch flangeability is lowered.

  Production No. No. 57 is an example in which since the amount of Ti is large, the ratio of non-recrystallized ferrite is high and sufficient ductility cannot be obtained. Production No. 58 is an example in which (Equation 2) is not satisfied, the strength of the hard tissue is not sufficient, and a satisfactory strength-ductility balance cannot be obtained.

  Production No. Nos. 59 and 60 are examples using steels that have a large total addition amount of Ti, Nb, Mo, and V and do not satisfy (Equation 1). In either case, the ratio of non-recrystallized ferrite is high, and sufficient ductility cannot be obtained. Production No. 61 is an example in which the amount of C is large, the proportion of the structure other than ferrite is large, and sufficient ductility and stretch flangeability cannot be obtained.

  Production No. 7, 10, 29 and 44 are comparative examples in which the conditions of the hot rolling process are outside the scope of the present invention. Production No. 44 is an example where the heating temperature of the hot rolling is low. Production No. 7 is an example where the finishing temperature is low. Production No. No. 29 has a slow cooling rate from the completion of rolling to 550 ° C. or less. 10 is an example in which the winding temperature is higher. In any of these examples, the Ti-based carbonitride becomes coarse, precipitation strengthening does not work on ferrite, and sufficient stretch flangeability is not obtained.

  Production No. No. 13 is a comparative example in which the cold rolling reduction ratio is insufficient, and the ratio of non-recrystallized ferrite is high and sufficient ductility cannot be obtained. Production No. No. 23 is an example in which since the annealing temperature is low, a hard structure is not obtained, and a large number of coarse cementite is generated, so that the hole expandability is lowered. Production No. 36 is an example in which the annealing temperature is high, the hard structure is increased, and the hole expandability is lowered. Production No. No. 26 is an example in which the cooling rate from 650 ° C. to 550 ° C. after annealing is low, pearlite increases, and the hole expandability decreases.

  The steel pieces shown in Table 1 and Table 4 were reheated, hot-rolled, cold-rolled, continuously annealed and plated under the conditions shown in Table 9 to obtain cold-rolled steel sheets. In Table 9, plating indicates the type of plating applied. Table 10 shows the microstructure observation results and materials of the steel sheet, and was evaluated in the same manner as in Examples 1 and 2.

  Production No. 62, 65 and 67 are examples in which hot dip galvanization is performed. Production No. 63, 64, and 66 are examples in which alloying hot dip galvanization was performed. As is apparent from Table 8, if the steel composition, hot rolling, cold rolling and annealing conditions are within the scope of the present invention, the strength-ductility balance and further the stretch flangeability can be achieved by plating after annealing. An excellent high-strength cold-rolled steel sheet can be obtained.

Claims (10)

% By mass
C: 0.05-0.20%,
Mn: 0.50 to 3.00%,
Ti: 0.03-0.15%,
Containing
Si: 2.50% or less,
Al: 1.50% or less,
P: 0.15% or less,
S: 0.010% or less,
N: 0.0060% or less,
Nb: 0.03% or less (including 0),
Mo: 0.25% or less (including 0),
V: 0.25% or less (including 0)
The content of C, N, Ti, Nb, Mo, V satisfies the following (formula 1) and the following (formula 2), the balance is composed of Fe and inevitable impurities, The particle diameter is 1 to 50 nm, the area ratio of ferrite is 50% or more, the area ratio of the hard structure composed of one or both of martensite and bainite is 5 to 50%, the remaining pearlite, residual austenite and The total area ratio of cementite is limited to 5% or less, the average grain size of the ferrite is limited to 20 μm or less, and the proportion of the unrecrystallized ferrite in the ferrite is limited to 25% or less. Precipitation strengthened double-phase cold-rolled steel sheet.
0.18 ≦ 6Ti + 25Nb + 3Mo + 3V ≦ 1.0 (Formula 1)
20C + 17.1N-5Ti-2.6Nb-2.5Mo-4.7V ≧ 0.6
... (Formula 2)
Here, C, N, Ti, Nb, Mo, and V are contents [mass%] of each element. % By mass
B: 0.0050% or less,
Cr: 2.00% or less,
Ni: 1.00% or less,
Cu: 1.00% or less,
The precipitation-strengthened double-phase cold-rolled steel sheet according to claim 1, characterized by containing one or more of W: 0.20% or less. % By mass
Ca: 0.010% or less,
One or both of REM: 0.100% or less is contained, The precipitation strengthening type | mold double phase cold rolled steel sheet of Claim 1 or 2 characterized by the above-mentioned. The precipitation strengthened double-phase cold-rolled steel sheet according to any one of claims 1 to 3, wherein the following (Formula 3) is satisfied.
80- (230C + 20Si + 55Mn-65Al + 125Ti
+ 1760Nb + 480Mo + 210V + 80000B + 35Cr
+ 25Ni + 20Cu + 25W) <0 (Formula 3)
Here, C, Si, Mn, Al, Ti, Nb, Mo, V, B, Cr, Ni, Cu, and W are contents [mass%] of each element.   A precipitation strengthened double-phase cold-rolled steel sheet, wherein the cold-rolled steel sheet according to any one of claims 1 to 4 is hot-dip galvanized.   A precipitation-strengthened double-phase cold-rolled steel sheet, wherein the cold-rolled steel sheet according to any one of claims 1 to 4 is subjected to alloying hot-dip galvanizing. It is a manufacturing method of the cold rolled steel plate of any one of Claims 1-4, Comprising: The steel plate which has a chemical component of any one of Claims 1-4 is heated to 1050 degreeC or more, The steel sheet is cooled to 550 ° C. or lower at a cooling rate of 15 ° C./s or higher, and is further subjected to cold rolling with a total rolling reduction of 40% or higher, and the maximum heating temperature exceeds (Ac ₁ +20) ° C. to (Ac ₃ − 20) A method for producing a precipitation-strengthened double-phase cold-rolled steel sheet, characterized by performing continuous annealing at a temperature of 650 ° C. to 500 ° C. and an average cooling rate of 3 ° C./s or higher. It is a manufacturing method of the cold rolled steel plate of any one of Claims 1-4, Comprising: The steel slab which has the chemical component of any one of Claims 1-4 is heated to 1050 degreeC or more. The steel sheet is hot-rolled at a finishing temperature of (Ae ₃ -30) ° C. or higher, cooled to 550 ° C. or lower at a cooling rate of 15 ° C./s or higher, and further cooled with a total rolling reduction of 40% or higher. Performing a continuous rolling, and performing a continuous annealing in which the maximum heating temperature is (Ac ₁ +20) ° C. to (Ac ₃ −20) ° C., and the average cooling rate from 650 ° C. to 500 ° C. is 3 ° C./s or more. A method for producing a precipitation-strengthened double-phase cold-rolled steel sheet.   It is a manufacturing method of the cold rolled steel sheet of Claim 5, Comprising: The cold rolled steel sheet manufactured by the manufacturing method of Claim 7 or 8 is immersed in a hot dip galvanizing bath after continuous annealing. A method for producing a precipitation-strengthened double-phase cold-rolled steel sheet.   A method for producing a cold-rolled steel sheet according to claim 6, wherein the precipitation-strengthened double-phase cold-rolled steel sheet is subjected to an alloying treatment after being immersed in the hot dip galvanizing bath according to claim 9. Production method.

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