JP2015147958A - High proportional limit high-strength cold-rolled thin steel sheet and production method thereof - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high proportional limit high-strength cold-rolled thin steel sheet and a production method thereof.SOLUTION: The high proportional limit high-strength cold-rolled thin steel sheet has: a composition in which, in mass%, C: 0.04 to 0.25%, Si: 0.30% or less, Mn: 0.1 to 2.0%, Al: 0.10% or less, and N: 0.010% or less, are contained, and furthermore one or two selected from Ti: 0.01 to 1.00%, and V: 0.01 to 1.00%, are contained such that (12/48)×Ti+(12/51)×V≥0.04 and C≥0.9×((12/48)×Ti+(12/51)×V) are satisfied; and a structure in which the ferrite phase occupying 95% or more in an area ratio, is the main phase, the deposition of less than 10 nm is 5.0×10μmor more, and the dislocation density is 5.0×10mor more. Thereby, a high proportional limit cold-rolled thin steel sheet having tensile strength: 980 MPa or more and a proportional limit of 0.6TS or more is obtained.

Description

  The present invention relates to a high-strength cold-rolled thin steel sheet, in particular, frame members such as automobile pillars and side sills, members and their reinforcing members, automobile door impact beams, and vending machines, desks, home appliances / OA devices, The present invention relates to a high-proportional high-strength cold-rolled thin steel sheet suitable for structural members used for building materials and the like, and a method for producing the same.

In recent years, in response to growing interest in the preservation of the global environment, there is a strong desire to reduce the amount of steel sheets that emit a large amount of carbon dioxide (CO ₂ ) during production. Further, in the automobile field and the like, there is an increasing demand for reducing the weight of the vehicle body, improving fuel efficiency, and reducing the amount of exhaust gas.
In response to such demands, attempts are being made to reduce the weight of automobile bodies by reducing the thickness of the steel sheet by applying a high-strength steel sheet. In particular, when a high-strength steel plate is applied to a structural member (component), the strength of the member increases, so that there is an advantage that collision absorption energy can be increased or plastic deformation can be suppressed. ) There is a problem that rigidity decreases as the wall thickness decreases. In particular, when a load is applied exceeding the proportional limit, the deflection of the member (part) increases. Therefore, it can be said that it is effective to increase the proportional limit of the steel sheet from the viewpoint of preventing the decrease in rigidity.

  As a steel plate with a high proportional limit, for example, in Patent Document 1, in mass%, C: 0.001 to 0.20%, Si: 1.5% or less, Mn: 2.5% or less, P: 0.010% or less, S: 0.01% or less , Acid-soluble Al: 0.001 to 0.10%, N: Hot-rolled steel sheet with a composition containing 0.015% or less, pickled, cold-rolled at 10 to 80%, and then cold-rolled at 200 to 600 ° C The manufacturing method of the high proportionality limit steel plate excellent in the bending workability heated to the temperature of this range is described. In the technique described in Patent Document 1, when cold rolling is performed on a steel sheet having a specific composition range, and annealing is performed in a relatively low temperature region below the recrystallization temperature, solute C and solute are dissolved simultaneously with rearrangement of dislocations. Movable dislocations are fixed due to the precipitation of N, and the bending limit is improved at the same time as the proportional limit is improved.

Patent Document 2 discloses that hot rolled steel sheets containing C: 0.01% and 0.05% or less by mass are cold-rolled, and then at a temperature of 700 ° C. to Ac ₃ transformation point for 20 to 120 seconds. A method for producing a low-carbon steel sheet excellent in formability and tensile rigidity is described in which recrystallization annealing is performed and overaging treatment is further performed in a temperature range of 250 to 500 ° C. for 1 to 5 minutes. According to the technique described in Patent Document 2, the yield strength is 120 to 250 N / mm ² , 2% pre-strain is applied at an equivalent strain, and heat treatment is performed at 150 to 170 ° C. for 5 to 20 minutes. Later, the amount of strain when the tensile test is performed again: The ratio of the slope X of the stress strain curve at 0.06% to the Young's modulus Y, X / Y (= A) exceeds 0.8, and the yield stress increases due to heat treatment A method for producing a low-carbon steel sheet excellent in formability and tensile rigidity having a cost of 40 N / mm ² or more is described. According to the technique described in Patent Document 2, it is possible to suppress the decrease in instantaneous Young's modulus and improve the tension rigidity by promoting segregation of interstitial solid solution elements around movable dislocations by heat treatment after molding. Yes.

JP 2010-138444 A JP 2001-348644

However, the techniques described in Patent Document 1 and Patent Document 2 use solid solution C or solid solution N to increase the proportional limit, but the amount of solid solution of C or N in steel is limited. Therefore, there is a limit to further strengthening.
An object of the present invention is to solve the above-described problems of the prior art and to provide a high-strength cold-rolled thin steel sheet having a high proportional limit and a method for manufacturing the same.

  Here, “high strength” refers to the case where the tensile strength TS is 980 MPa or more, and “thin steel plate” refers to the case where the thickness is 2.0 mm or less. The cases where (proportional limit / tensile strength) is 0.6 or more shall be referred to.

  In order to solve the above-described problems, the present inventors diligently studied various factors that affect the proportional limit in the cold-rolled steel sheet. As a result, it has been found that movable dislocations can be fixed by fine precipitates of less than 10 nm, thereby suppressing plastic deformation during stress loading and dramatically increasing the proportional limit. Specifically, it is possible to dramatically increase the proportional limit by depositing a large amount of fine precipitates containing Ti and / or V and less than 10 nm and increasing the dislocation density.

FIG. 1 shows the relationship between the proportional limit / TS and the precipitation density of fine precipitates of less than 10 nm. FIG. 1 shows the case where the dislocation density is 5.0 × 10 ¹⁴ / m ² or more. FIG. 1 shows that the proportional limit / TS becomes 0.6 or more by dispersing and precipitating fine precipitates of less than 10 nm in an amount of 5.0 × 10 ⁴ pieces / μm ³ or more. FIG. 2 shows the relationship between the proportional limit / TS and the dislocation density. FIG. 2 shows the case where the fine precipitates are 5.0 × 10 ⁴ pieces / μm ³ or more. FIG. 2 shows that the proportional limit / TS is 0.6 or more when the dislocation density is 5.0 × 10 ¹⁴ / m ² or more.

The present invention has been completed based on the above findings and further studies. That is, the gist of the present invention is as follows.
(1) Cold-rolled thin steel sheet in mass%, C: 0.04 to 0.25%, Si: 0.30% or less, Mn: 0.1 to 2.0%, P: 0.05% or less, S: 0.030% or less, Al: 0.10 %, N: 0.010% or less, Ti: 0.01 to 1.00%, V: 0.01 to 1.00% selected from one or two, and C, Ti, V Formula (1) and the following formula (2)
(12/48) × Ti + (12/51) × V ≧ 0.04 (1)
C ≧ 0.9 × ((12/48) × Ti + (12/51) × V) (2)
(Here, C, Ti, V: Content of each element (mass%))
And a composition comprising the balance Fe and unavoidable impurities, and a ferrite phase with an area ratio of 95% or more as a main phase, and the main phase and the second phase with an area ratio of 0 to 5% And a structure having a precipitate density of less than 10 nm of 5.0 × 10 ⁴ μm ⁻³ or more and a dislocation density of 5.0 × 10 ¹⁴ m ⁻² or more, and a tensile strength of 980 MPa or more. A high-strength cold-rolled steel sheet characterized by a high proportional limit.

(2) In (1), in addition to the above composition, it is further selected by mass% from Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005-0.600% A high-strength cold-rolled thin steel sheet characterized by containing one or more kinds.
(3) The high-strength cold-rolled thin steel sheet according to (1) or (2), further containing B: 0.0002 to 0.0050% by mass% in addition to the above composition.

(4) In any one of (1) to (3), in addition to the above composition, it is further selected by mass: Cr: 0.01-1.0%, Ni: 0.01-1.0%, Cu: 0.01-1.0% A high-strength cold-rolled thin steel sheet characterized by containing one or more kinds.
(5) The high strength cold-rolled thin steel sheet according to any one of (1) to (4), further containing Sb: 0.005 to 0.050% by mass in addition to the above composition.

(6) In any one of (1) to (5), one or two selected from Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01%, in addition to the above composition, in mass% A high-strength cold-rolled thin steel sheet characterized by containing.
(7) A high-strength cold-rolled thin steel sheet, wherein the high-strength cold-rolled thin steel sheet according to any one of (1) to (6) is a cold-rolled annealed sheet that has been subjected to a continuous annealing treatment.

(8) The high-strength cold-rolled thin steel sheet according to any one of (1) to (7), wherein a plating layer is formed on the steel sheet surface.
(9) When hot rolling and cold rolling are performed on the steel material to form a cold-rolled thin steel sheet, the steel material is mass%, C: 0.04 to 0.25%, Si: 0.30% or less, Mn: 0.1 to 2.0%, P: 0.05% or less, S: 0.030% or less, Al: 0.10% or less, N: 0.010% or less, further selected from Ti: 0.01-1.00%, V: 0.01-1.00% 1 or 2 types, and C, Ti and V are represented by the following formulas (1) and (2)
(12/48) × Ti + (12/51) × V ≧ 0.04 (1)
C ≧ 0.9 × ((12/48) × Ti + (12/51) × V) (2)
(Here, C, Ti, V: Content of each element (mass%))
And a slab having a composition comprising the balance Fe and inevitable impurities, and the hot rolling is a rolling consisting of rough rolling and finish rolling, and the finish rolling finish finish rolling temperature is 850 ° C. The rolling is as described above, and is cooled at a cooling rate of 30 ° C./s or higher at an average cooling rate from the finish rolling to 700 ° C., and is taken up at a winding temperature of 500 ° C. or more. , After subjecting the hot-rolled sheet to pickling treatment, cold rolling at a cold pressure ratio of 10 to 80% to obtain a cold-rolled sheet, tensile strength: 980 MPa or more, A method for producing high-strength cold-rolled steel sheets with a high proportional limit.

(10) In (9), the cold-rolled sheet is further subjected to a continuous annealing treatment with a soaking temperature of 600 to 800 ° C. and a soaking time of 300 s or less to obtain a cold-rolled annealed sheet. A manufacturing method of high strength cold-rolled thin steel sheet.
(11) In (10), during the cooling of the continuous annealing treatment, the steel plate is immersed in a galvanizing bath having a bath temperature of 420 to 500 ° C. and subjected to a hot dip galvanizing treatment to form a hot dip galvanized layer on the surface of the steel plate. A method for producing a high-strength cold-rolled thin steel sheet, characterized by using a steel sheet.

  (12) The method for producing a high-strength cold-rolled thin steel sheet according to (11), wherein after the hot-dip galvanizing treatment, an alloying treatment of the plating layer is further performed.

  According to the present invention, a high strength cold-rolled steel sheet having a tensile strength TS: 980 MPa or higher and a proportional limit as high as 0.6 TS or higher can be manufactured stably and at low cost, which is remarkably industrially superior. There is an effect. According to the present invention, it is possible to expect an effect that a decrease in rigidity due to thinning can be suppressed.

It is a graph which shows the influence of the precipitation density which has on proportionality limit / TS. It is a graph which shows the influence of the dislocation density which acts on proportional limit / TS.

The high-strength cold-rolled thin steel sheet of the present invention is, in mass%, C: 0.04 to 0.25%, Si: 0.30% or less, Mn: 0.1 to 2.0%, P: 0.05% or less, S: 0.030% or less, Al: 0.10% Hereinafter, it contains N: 0.010% or less, and further contains one or two selected from Ti: 0.01-1.00%, V: 0.01-1.00%, and C, Ti, and V are the following ( 1) and next (2)
(12/48) × Ti + (12/51) × V ≧ 0.04 (1)
C ≧ 0.9 × ((12/48) × Ti + (12/51) × V) (2)
(Here, C, Ti, V: Content of each element (mass%))
And having a composition composed of the remaining Fe and unavoidable impurities.

First, the reason for limiting the composition of the high-strength cold-rolled thin steel sheet of the present invention will be described. Hereinafter, unless otherwise specified, mass% is simply expressed as%.
C: 0.04-0.25%
C is an element that combines with carbide-type elements such as Ti and V to form fine carbides and contributes to high strength. C also has the effect of lowering the ferrite transformation start temperature in cooling after hot rolling, thereby lowering the precipitation temperature of carbides and contributing to the refinement of carbides. The formed fine carbides fix the movable dislocations and contribute to the improvement of the proportional limit. In order to obtain such an effect, a content of 0.04% or more is required. In addition, Preferably it is 0.08% or more, More preferably, it is 0.15% or more. On the other hand, a large content exceeding 0.25% suppresses ferrite transformation and promotes transformation to bainite and martensite. For this reason, fine carbide formation with Ti, V, etc. is controlled. Moreover, a large amount of C content also deteriorates weldability. For this reason, C was limited to 0.25% or less. In addition, Preferably it is 0.20% or less. For these reasons, C is limited to the range of 0.04 to 0.25%.

Si: 0.30% or less
Si has a function of increasing the strength of the steel sheet by solid solution. In order to obtain such an effect, it is desirable to contain 0.005% or more which is higher than the impurity level. Further, Si is a ferrite-forming element, and when contained in a large amount, Si has an adverse effect of promoting ferrite transformation in cooling after hot rolling, increasing the precipitation temperature of carbides, and precipitating carbides coarsely. Furthermore, when Si is contained in a large amount, in the annealing after hot rolling, Si oxide is generated on the surface of the steel sheet, and an unplated portion is generated at the time of plating treatment. Such an adverse effect is acceptable if it is 0.30% or less. In addition, Preferably it is 0.1% or less, More preferably, it is 0.05% or less, More preferably, it is 0.03% or less.

Mn: 0.1-2.0%
Mn is an element having a function of increasing the strength of a steel sheet by dissolving in steel and further detoxifying harmful S as MnS by combining with S. In addition, Mn has an effect of contributing to refinement of carbide by lowering the ferrite transformation start temperature and lowering the precipitation temperature of carbide in cooling after rolling. In order to obtain such an effect, Mn needs to be contained by 0.1% or more. In addition, Preferably it is 0.3% or more. On the other hand, a large content exceeding 2.0% suppresses ferrite transformation and promotes transformation to bainite and martensite. For this reason, fine carbide formation with Ti, V, etc. is controlled. For this reason, Mn was limited to 2.0% or less. In addition, Preferably it is 1.5% or less, More preferably, it is 1.0% or less. For these reasons, Mn is limited to a range of 0.1 to 2.0%.

P: 0.05% or less
P is an element that segregates at the grain boundaries, deteriorates ductility and toughness, and adversely affects the steel sheet characteristics. P promotes ferrite transformation in cooling after rolling, raises the precipitation temperature of carbides, coarsens ferrite grains, and coarsely precipitates carbides. Moreover, P reduces weldability. For this reason, it is desirable to reduce as much as possible in the present invention, but such an adverse effect is acceptable up to 0.05%. For this reason, P was limited to 0.05% or less. In addition, Preferably it is 0.03% or less, More preferably, it is 0.01% or less.

S: 0.030% or less
S is an element that significantly reduces hot ductility, induces hot cracking, and significantly deteriorates surface properties. Furthermore, S not only contributes to the strength, but also forms coarse sulfides, thereby reducing the ductility and stretch flangeability, and further reducing the weldability. It is desirable to reduce. In addition, since such a bad influence becomes remarkable when S exceeds 0.030%, S was limited to 0.030% or less. In addition, Preferably it is 0.010% or less, More preferably, it is 0.003% or less, More preferably, it is 0.001% or less.

Al: 0.10% or less
Al is an element that acts as a deoxidizer, and as Al killed steel, it is desirable to contain 0.01% or more in order to obtain such an effect. In addition, Al has the effect of promoting ferrite transformation by cooling after rolling, thereby adversely affecting the coarsening of ferrite grains and the coarse precipitation of carbides through the increase in carbide precipitation temperature. . Therefore, it is necessary to avoid containing a large amount. Further, a large content exceeding 0.10% leads to an increase in aluminum oxide in the steel, and adverse effects such as a decrease in cleanliness and a decrease in ductility. For these reasons, Al is limited to 0.10% or less. In addition, Preferably it is 0.06% or less.

N: 0.010% or less
N forms coarse nitrides at high temperatures with Ti, V, etc., and contributes less to strength, while reducing contributions to higher strength due to Ti, V, etc. Furthermore, a large amount of N may induce slab cracking during hot rolling and generate surface defects. For these reasons, it is desirable to reduce N as much as possible, but it is acceptable up to 0.010%. For this reason, N was limited to 0.010% or less. In addition, Preferably it is 0.005% or less, More preferably, it is 0.003% or less, More preferably, it is 0.002% or less.

One or two selected from Ti: 0.01-1.00%, V: 0.01-1.00%
Ti and V are elements that combine with C to form fine carbides and contribute to increasing strength and increasing the proportional limit. In the present invention, either Ti or V is contained. . In order to obtain such an effect, it is necessary to contain Ti: 0.01% or more and V: 0.01% or more. On the other hand, even if contained in a large amount exceeding Ti: 1.00%, V: 1.00%, the increase in strength and proportional limit is saturated, and an effect commensurate with the content cannot be expected, which is economically disadvantageous. For this reason, it was limited to Ti: 1.00% or less and V: 1.00% or less. For these reasons, Ti is limited to the range of 0.01 to 1.00% and V: 0.01 to 1.00%.

C, Ti, and V are contained within the above range and adjusted so as to satisfy the following formulas (1) and (2).
(12/48) x Ti + (12/51) x V≥ 0.04 (1)
C ≧ 0.9 × ((12/48) × Ti + (12/51) × V) (2)
(Here, C, Ti, V: Content of each element (mass%))
The left side {(12/48) × Ti + (12/51) × V} of the formula (1) is a value obtained by converting the Ti and V contents into the amount of carbon that forms carbides in terms of atomic ratio, and the converted amount of carbon. It is. If this carbon amount conversion value is less than 0.04, the Ti and V contents are too small to ensure the desired high strength and the desired proportional limit. For this reason, {(12/48) × Ti + (12/51) × V} is limited to 0.04 or more. In addition, Preferably it is 0.06 or more, More preferably, it is 0.10 or more. On the other hand, when {(12/48) × Ti + (12/51) × V} exceeds 0.20, the increase in strength and proportional limit does not increase so much, and an effect commensurate with the content cannot be expected. For this reason, {(12/48) × Ti + (12/51) × V} is preferably 0.20 or less. In addition, More preferably, it is 0.15 or less.

  Furthermore, if the Ti and V contents are higher than the C content by atomic ratio, the proportion of Ti and V contained will precipitate as carbides, the resulting precipitates will become coarse, and the desired high strength and proportional limit will be achieved. It cannot be secured. Therefore, the C content is 0.9 times or more of {(12/48) × Ti + (12/51) × V}, which is the carbon amount converted value of Ti and V, so as to satisfy the formula (2). It was decided to adjust. In addition, Preferably it is 1.0 times or more, More preferably, it is 1.1 times or more.

For this reason, the contents of C, Ti, and V were adjusted so as to satisfy the expressions (1) and (2) within the above-described content range.
The above components are basic components. In the present invention, Nb: 0.005 to 0.600%, Mo: 0.005 to 0.600%, Ta: 0.005 to 0.600%, W: One or more selected from 0.005 to 0.600% and / or B: 0.0002 to 0.0050% and / or Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cu: 0.01 to 1 or more selected from 1.0%, and / or Sb: 0.005 to 0.050%, and / or Ca: 0.0005 to 0.01%, REM: 1 selected from 0.0005 to 0.01% Species or two can be contained.

One or more selected from Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005-0.600%
Nb, Mo, Ta, and W are all elements that form fine precipitates and contribute to the increase in strength and the increase in proportionality by precipitation strengthening. Select one or more as required. Can be contained. In order to obtain such an effect, it is preferable to contain Nb: 0.005% or more, Mo: 0.005% or more, Ta: 0.005% or more, and W: 0.005% or more. More preferably, Nb is 0.100% or more, Mo is 0.100% or more, Ta is 0.100% or more, and W is 0.100% or more. On the other hand, even if it contains Nb: 0.600%, Mo: 0.600%, Ta: 0.600%, W: 0.600% in large amounts, the effect is saturated and an effect commensurate with the content cannot be expected. Disadvantageous. For this reason, when it contains, it is preferable to limit to Nb: 0.600% or less, Mo: 0.600% or less, Ta: 0.600% or less, W: 0.600% or less, respectively. More preferably, Nb is 0.300% or less, Mo is 0.300% or less, Ta is 0.300% or less, and W is 0.300% or less. Therefore, when it is contained, it is preferable to limit it to the range of Nb: 0.005 to 0.600%, Mo: 0.005 to 0.600%, Ta: 0.005 to 0.600%, W: 0.005 to 0.600%.

B: 0.0002 to 0.0050%
B has the effect of lowering the ferrite transformation start temperature in cooling after rolling, and contributes to the refinement of the carbide by lowering the precipitation temperature of the carbide. Further, B segregates at the grain boundary to increase the grain boundary strength, and contributes to improvement of secondary work embrittlement resistance. In order to obtain such an effect, B is preferably contained in an amount of 0.0002% or more. More preferably, it is 0.0005% or more, More preferably, it is 0.0010% or more. On the other hand, a large content exceeding 0.0050% increases hot deformation resistance, makes rolling difficult, and causes a decrease in ductility. For this reason, B is preferably limited to 0.0050% or less. In addition, More preferably, it is 0.0030% or less, More preferably, it is 0.0020% or less. Therefore, when contained, B is preferably limited to a range of 0.0002 to 0.0050%.

One or more selected from Cr: 0.01-1.0%, Ni: 0.01-1.0%, Cu: 0.01-1.0%
Cr, Ni, and Cu are all elements that contribute to increasing the strength through the refinement of the structure, and can be selected as necessary to contain one or more. In order to obtain such an effect, each content is preferably 0.01% or more. On the other hand, even if contained in a large amount exceeding 1.0%, the effect is saturated, and an effect commensurate with the content cannot be expected, which is economically disadvantageous. For this reason, when it is contained, it is preferable to limit to Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, and Cu: 0.01 to 1.0%, respectively.

Sb: 0.005 to 0.050%
Sb segregates on the surface of the slab (steel material) during hot rolling, prevents slab nitriding, and suppresses the formation of coarse nitrides, and can be contained as necessary. In order to acquire such an effect, it is preferable to contain 0.005% or more. On the other hand, if the content exceeds 0.050%, the material cost increases. For these reasons, when contained, Sb is preferably limited to a range of 0.005 to 0.050%.

One or two selected from Ca: 0.0005-0.01%, REM: 0.0005-0.01%
Both Ca and REM are elements that contribute to the improvement of ductility and stretch flangeability through the form control of sulfides, and can be selected and contained as necessary. In order to acquire such an effect, it is preferable to contain Ca: 0.0005% or more and REM: 0.0005% or more, respectively. On the other hand, even if each content exceeds 0.01%, the effect is saturated, and an effect commensurate with the content cannot be expected, which is economically disadvantageous. For this reason, when it contains, it is preferable to limit to Ca: 0.0005-0.01% and REM: 0.0005-0.01%, respectively.

The balance other than the components described above consists of Fe and inevitable impurities. Examples of unavoidable impurities include Sn, Mg, Co, As, Pb, Zn, O, etc. However, these impurity elements are acceptable as long as the total is 0.5% or less, with no problem in characteristics.
The high-strength cold-rolled steel sheet of the present invention has the above-described composition, and further has a ferrite phase with an area ratio of 95% or more as a main phase, and the main phase and the second phase with an area ratio of 0 to 5% It has an organization consisting of

Main phase: ferrite phase having an area ratio of 95% or more In the present invention, a ferrite phase having an area ratio of 95% or more is used as a main phase in order to improve moldability. If a low-temperature transformation phase such as a bainite phase or a martensite phase is present as the second phase, movable dislocations are introduced during transformation and the proportional limit is lowered. For this reason, in order to increase the proportional limit, the area ratio of the second phase other than the ferrite phase is preferably 0%. Even if it exists, the area ratio is 5% or less. For this reason, the ferrite phase as the main phase is set to 95% or more in area ratio. The area ratio is preferably 98% or more, and more preferably 100%.

The high-strength cold-rolled thin steel sheet of the present invention has a structure in which the precipitate density of less than 10 nm is 5.0 × 10 ⁴ μm ⁻³ or more and the dislocation density is 5.0 × 10 ¹⁴ m ⁻² or more.
Particle size: Precipitation density of precipitates of less than 10 nm: 5.0 × 10 ⁴ μm ⁻³ or more In the present invention, movable dislocations are fixed with fine precipitates to increase the proportional limit. Coarse precipitates have little impact on strength. In the present invention, the fine precipitate is a precipitate having a particle size of less than 10 nm. The preferred particle size is less than 5 nm, more preferably less than 3 nm. The “particle diameter” here is the maximum diameter. Particle size: If the precipitation density of precipitates of less than 10 nm is less than 5.0 × 10 ⁴ μm ⁻³ , the desired high proportional limit cannot be secured. For this reason, the precipitation density of precipitates having a particle size of less than 10 nm was limited to 5.0 × 10 ⁴ μm ⁻³ or more. In addition, it is preferably 10.0 × 10 ⁴ μm ⁻³ or more, more preferably 20.0 × 10 ⁴ μm ⁻³ or more.

Dislocation density: 5.0 × 10 ¹⁴ m ⁻² or more In the present invention, as described above, dislocations introduced by processing such as cold rolling are fixed with fine precipitates having a particle size of less than 10 nm to obtain a desired high proportional limit. . In order to secure a desired proportional limit, a dislocation density of 5.0 × 10 ¹⁴ m ⁻² or more is required. If the density of the introduced dislocations is less than 5.0 × 10 ¹⁴ m ⁻² , the proportional limit is hardly increased and the desired proportional limit cannot be secured. For this reason, the dislocation density is limited to 5.0 × 10 ¹⁴ m ⁻² or more. In addition, it is preferably 10.0 × 10 ¹⁴ m ⁻² or more, more preferably 20.0 × 10 ¹⁴ m ⁻² or more, and further preferably 30.0 × 10 ¹⁴ m ⁻² or more. Also, even if dislocations exceeding 100.0 × 10 ¹⁴ m ⁻² are introduced, the effect is saturated and an effect commensurate with the dislocation introduction process cannot be expected. Therefore, a dislocation density of 100.0 × 10 ¹⁴ m ⁻² or less is sufficient. It is. The dislocations present in the cold-rolled thin steel sheet of the present invention are dislocations that are not tangled and remain random.

  In the high strength cold-rolled thin steel sheet of the present invention, a plating layer may be formed on the steel sheet surface. As the plating layer, a hot dip galvanized layer, an alloyed hot dip galvanized layer, an electrogalvanized layer and the like are all suitable. Needless to say, a film such as a chemical conversion treatment may be formed. The plating layer to be formed may be a zinc / Al composite plating layer, a zinc / Ni composite plating layer, an Al plating layer, an Al / Si composite plating layer, or the like.

Next, production conditions for the high-strength cold-rolled steel sheet of the present invention will be described.
In the present invention, the steel material having the above composition is subjected to hot rolling and cold rolling to obtain a cold rolled thin steel sheet.
The method for producing the steel material is not particularly limited, and any conventional method can be applied.
For example, it is possible to exemplify a conventional method in which molten steel having the above composition is melted by a melting method using a converter or the like, and a slab or the like is cast by a conventional casting method such as a continuous casting method to obtain a steel material. .

The obtained steel material is hot rolled.
When the obtained steel material maintains a temperature at which hot rolling is possible, hot rolling may be performed as it is. In addition, when the steel material is a hot piece or a cold piece, the steel material is reheated to a predetermined heating temperature and then hot rolled.
The heating temperature for hot rolling is preferably 1200 ° C. or higher and lower than the melting point. By setting the heating temperature to a high temperature, the carbide-forming element can be completely dissolved, and can be precipitated as fine precipitates in the subsequent process. If the heating temperature is less than 1200 ° C, the carbide-forming element cannot be completely dissolved. The heating of the steel material is preferably maintained at 1200 ° C or higher for 10 minutes or more, more preferably 1250 ° C or higher for 30 minutes or longer, more preferably 1300 ° C or higher for 10 minutes or longer, more preferably 1400 ° C or higher. 10min or more.

The heated steel material is then subjected to hot rolling comprising rough rolling and finish rolling to form a hot rolled sheet.
The rough rolling is not particularly limited, and any conventional rough rolling can be applied.
After rough rolling to obtain a sheet bar having a predetermined size, finish rolling is performed to obtain a hot rolled sheet. In the finish rolling, the finish rolling finish temperature is 850 ° C. or higher.

Finish rolling end temperature: 850 ° C or higher When the finish rolling end temperature is lower than 850 ° C, ferrite transformation is promoted by cooling after rolling, resulting in coarse carbide with accelerated carbide precipitation, and the desired high proportional limit is achieved. Can't be achieved. In addition, when the finish rolling finish temperature is low enough to be in the ferrite region, coarse carbides precipitate due to strain-induced precipitation. For this reason, the finish rolling end temperature is limited to 850 ° C. or higher. In addition, Preferably it is 880 degreeC or more, More preferably, it is 920 degreeC or more, More preferably, it is 940 degreeC or more.

After finishing rolling, the temperature range up to 700 ° C. is cooled at an average cooling rate of 30 ° C./s or more, and the coil is wound into a coil at a winding temperature of 500 ° C. or more to obtain a hot rolled sheet. In the present invention, fine precipitates are deposited in the winding process.
Average cooling rate from finish rolling to 700 ° C: 30 ° C / s or more When the cooling in the temperature range from finish rolling to 700 ° C is slower than 30 ° C / s on average, ferrite transformation is promoted Carbide is largely precipitated. For this reason, in the present invention, cooling after finishing rolling is limited to 30 ° C./s or more at an average cooling rate from finishing finishing to 700 ° C. In addition, Preferably it is 50 degrees C / s or more, More preferably, it is 70 degrees C / s or more. The upper limit of the cooling rate is not particularly limited, but is preferably 1000 ° C./s or less from the viewpoint of steel plate shape and temperature control.

Winding temperature: 500 ° C or higher If the winding temperature is lower than 500 ° C, the formation of low-temperature transformation phases such as bainite and martensite is promoted, and the formation of the desired structure with the ferrite phase as the main phase is hindered. Carbide precipitation is also suppressed. For this reason, the coiling temperature was limited to 500 ° C. or higher. In addition, Preferably it is 550 degreeC or more, More preferably, it is 600 degreeC or more. On the other hand, from the viewpoint of suppressing the coarsening of the carbide, the coiling temperature is preferably 700 ° C. or lower. More preferably, it is 650 ° C. or lower.

The obtained hot-rolled sheet is then subjected to pickling treatment and cold rolling at a cold pressure ratio of 10 to 80% to obtain a cold-rolled sheet. As the pickling treatment, any conventional pickling can be applied.
Cold pressure ratio: 10-80%
Cold rolling is performed to introduce dislocations into the steel sheet. With the introduction of dislocations, an increase in strength and proportional limit can be expected. In order to obtain such an effect, the rolling reduction (cold pressure ratio) of cold rolling is set to 10% or more. In addition, Preferably it is 20% or more, More preferably, it is 30% or more. If the cold pressure ratio is less than 10%, the desired high strength and high proportional limit cannot be obtained. On the other hand, if the cold pressure ratio exceeds 80%, the workability deteriorates remarkably. For this reason, the cold pressure rate was limited to 80% or less. In addition, Preferably it is 70% or less.

In addition, the obtained cold-rolled sheet may be further subjected to an annealing treatment to form a cold-rolled annealed sheet. Annealing treatment is performed to improve formability. The annealing treatment is continuous annealing from the viewpoint of productivity. The conditions for the continuous annealing treatment are preferably soaking temperature: 600 to 800 ° C., soaking time: 300 s or less.
Soaking temperature: 600-800 ° C
The soaking temperature in the annealing treatment is set to 600 ° C. or higher in order to ensure the desired formability. When the soaking temperature is less than 600 ° C., no recovery occurs, and thus no improvement in moldability is observed. On the other hand, when the soaking temperature is higher than 800 ° C., the precipitates become coarse and recrystallization progresses, so that the strength and the proportional limit decrease. For this reason, the soaking temperature of the annealing treatment was limited to a temperature in the range of 600 to 800 ° C. In addition, Preferably it is 780 degrees C or less, More preferably, it is 760 degrees C or less.

The heating rate up to the soaking temperature is not particularly limited, but is preferably 1 to 100 ° C./s from the viewpoint of productivity and temperature control.
Soaking time: 300 s or less In order to ensure the desired formability improvement, the soaking time is desirably 30 s or more. If the soaking time is less than 30 s, recovery does not occur, and the desired formability cannot be improved. On the other hand, for a long time exceeding 300 s, fine precipitates become coarse, recrystallization proceeds, and the dislocation density also decreases. For this reason, the soaking time was set to 30 to 300 s. In addition, Preferably it is 200 s or less, More preferably, it is 100 s or less.

The cooling rate from the soaking temperature is not particularly limited, but is preferably in the range of 1 to 100 ° C./s from the viewpoint of productivity and temperature control.
When hot dip galvanizing treatment is performed, the steel sheet is immersed in a galvanizing bath having a bath temperature of 420 to 500 ° C. during the cooling of the annealing treatment (continuous annealing treatment) to form a plating layer on the steel plate surface. It is preferable. In addition, it is preferable to perform the hot dip galvanization process using the continuous galvanization equipment which has arrange | positioned the galvanization bath to the exit side of the soaking zone.

Moreover, you may perform the alloying process of a hot dip galvanized layer as needed to the hot dip galvanized steel plate which hot-dip galvanized process was performed and the hot dip galvanized layer was formed in the surface.
Heating temperature for alloying treatment: 460-600 ° C
In order to alloy Zn and Fe, it is necessary to heat the plating layer to a temperature of 460 ° C. or higher. On the other hand, when the heating temperature is higher than 600 ° C., alloying proceeds too much and the plating becomes brittle. For this reason, the alloying treatment of the hot dip galvanized layer is performed in the temperature range of 460 to 600 ° C. In addition, Preferably it is 570 degrees C or less. Holding at the above-described temperature is 1 s or more, preferably 10 s or less. When the holding time at the heating temperature of the alloying treatment exceeds 10 s and becomes a long time, precipitates in the ground iron become coarse, and the strength of the steel sheet decreases. For this reason, it is preferable that the holding time in an alloying process shall be 1-10 s.

  Hereinafter, based on an Example, this invention is further demonstrated.

  Molten steel having the composition shown in Table 1 was melted in a converter, and a steel material made into a slab by a continuous casting method was used as a starting material. These steel materials were hot-rolled and cold-rolled under the conditions shown in Table 2 to obtain cold-rolled steel sheets having the thicknesses shown in Table 2. Needless to say, the hot-rolled sheet was pickled before cold rolling. Further, some of the steel sheets were further subjected to annealing treatment, further galvanizing treatment, or further alloying treatment of the plating layer under the conditions shown in Table 2.

Test pieces were collected from the obtained steel plates, subjected to structure observation and tensile tests, and evaluated for strength characteristics. The test method was as follows.
(1) Microstructure observation A test specimen for microstructural observation was collected from the obtained steel sheet, the cross section in the rolling direction was polished and subjected to nital corrosion, and the microstructure was observed using an optical microscope (magnification: 500 times). 300 × 300 μm ² The structure of the region was imaged, and the area ratio of ferrite was calculated.

Moreover, from the obtained steel plate, a thin film sample was prepared by grinding, mechanical polishing, electrolytic polishing, etc., and using a transmission electron microscope (magnification: 300,000 times), 10 regions of 100 × 100 nm ² area, Precipitates were observed and the number of precipitates less than 10 nm was measured. At that time, the film thickness in the measurement field of view was also obtained by a convergent electron diffraction method, and the precipitation density (pieces / μm ³ ) of precipitates of less than 10 nm was calculated. Further, the diameter of 500 precipitates less than 10 nm was measured, and the average particle diameter was obtained by arithmetic averaging. In measuring the particle size of the precipitate, since the precipitate was not spherical, the maximum value was measured and used as the particle size of the precipitate.

In addition, from the obtained steel plate, a specimen for X-ray diffraction (size: thickness x 20 mm x 20 mm) is collected by grinding, mechanical polishing, and chemical polishing so that the 1/4 thickness is the surface. The local strain η is obtained by the X-ray diffraction method using a Cu tube, and the following equation is given: ρ = 14.4 × η ² / b ²
Was used to measure the dislocation density ρ. Here, b is a Burgers vector (0.25 nm). As the local strain, a value obtained by using the Hall method (WHHall: J. Inst. Met., 77 (1950), 1127) is used. Moreover, the above-mentioned formula used the formula by GKWilliams and RESmallmann (GKWilliams and RESmallmann: Philos. Mag., 8 (1956), 34).
(2) Tensile test A tensile test piece (JIS No. 5 test piece) is taken from the obtained steel sheet so that the direction perpendicular to the rolling direction (C direction) is the longitudinal direction of the test piece, and stipulated in JIS Z 2241 In accordance with the tensile test, tensile properties (yield strength YP, tensile strength TS, elongation El) were determined. Moreover, from the obtained steel plate, a tensile test piece (JIS No. 5 test piece) was sampled so that the direction perpendicular to the rolling direction (C direction) was the test piece longitudinal direction, and the proportional limit was obtained. Strain gauges (gauge length: 5 mm) are affixed to both sides of the parallel part of a tensile test piece, a tensile test is performed at a strain rate of 1 mm / min, a stress-strain curve is obtained, and the slope of the obtained stress-strain curve is The stress at the point at which it began to decrease (the point at which it started to deviate from the straight line) was defined as proportional limit stress. The strain was an average value of values measured with strain gauges attached to both sides of the test piece.

  The obtained results are shown in Table 3.

  All of the inventive examples are high-strength cold-rolled thin steel sheets having a tensile strength TS: 980 MPa or more and a high proportional limit of (0.6 × TS) MPa or more. On the other hand, in the comparative example that is outside the scope of the present invention, the desired tensile strength is not obtained, or the desired high proportional limit is not ensured, or both the desired tensile strength and the desired proportional limit are obtained. It is not done.

Claims (12)

Cold rolled thin steel sheet in mass%,
C: 0.04-0.25%, Si: 0.30% or less,
Mn: 0.1 to 2.0%, P: 0.05% or less,
S: 0.030% or less, Al: 0.10% or less,
N: not more than 0.010%, further containing one or two selected from Ti: 0.01 to 1.00%, V: 0.01 to 1.00%, and C, Ti and V in the following (1) And a composition comprising the balance Fe and unavoidable impurities, adjusted to satisfy the formula and the following formula (2):
The ferrite phase with an area ratio of 95% or more is the main phase, and consists of the main phase and the second phase with an area ratio of 0 to 5%. Further, the precipitation density of precipitates of less than 10 nm is 5.0 × 10 ⁴ μm ⁻ Having a structure of ³ or more and a dislocation density of 5.0 × 10 ¹⁴ m ⁻² or more,
Tensile strength: 980 MPa or higher, high strength cold-rolled thin steel sheet characterized by high proportional limit.
Record
(12/48) × Ti + (12/51) × V ≧ 0.04 (1)
C ≧ 0.9 × ((12/48) × Ti + (12/51) × V) (2)
Here, C, Ti, V: Content of each element (mass%)   In addition to the above-mentioned composition, by mass%, Nb: 0.005 to 0.600%, Mo: 0.005 to 0.600%, Ta: 0.005 to 0.600%, W: 0.005 to 0.600% The high-strength cold-rolled thin steel sheet according to claim 1, which is contained. The high-strength cold-rolled steel sheet according to claim 1 or 2, further comprising B: 0.0002 to 0.0050% by mass% in addition to the composition. In addition to the above composition, it is further mass%, Cr: 0.01-1.0%, Ni: 0.01-1.0%, Cu: 0.01-1.0%
The high-strength cold-rolled steel sheet according to any one of claims 1 to 3, comprising one or more selected from among the above.   The high-strength cold-rolled thin steel sheet according to any one of claims 1 to 4, further comprising Sb: 0.005 to 0.050% in addition to the composition.   6. In addition to the above composition, the composition further contains one or two selected from Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01% by mass%. The high-strength cold-rolled thin steel sheet described in 1.   The high-strength cold-rolled thin steel sheet according to any one of claims 1 to 6, wherein the high-strength cold-rolled thin steel sheet is a cold-rolled annealed sheet that has been subjected to a continuous annealing treatment.   The high strength cold-rolled thin steel sheet according to any one of claims 1 to 7, wherein a plated layer is formed on the surface of the steel sheet. The steel material is subjected to hot rolling and cold rolling to form a cold rolled steel sheet.
The steel material in mass%,
C: 0.04-0.25%, Si: 0.30% or less,
Mn: 0.1 to 2.0%, P: 0.05% or less,
S: 0.030% or less, Al: 0.10% or less,
N: not more than 0.010%, further containing one or two selected from Ti: 0.01 to 1.00%, V: 0.01 to 1.00%, and C, Ti and V in the following (1) And a steel material having a composition comprising the balance Fe and inevitable impurities, adjusted to satisfy the formula and the following formula (2):
The hot rolling is rolling consisting of rough rolling and finish rolling, and the finish rolling is a finish rolling finish temperature: 850 ° C. or higher, and the average cooling rate from the finish rolling finish to 700 ° C. is 30 ° C./s. Cool at the above cooling rate, take-up temperature: 500 ° C or higher to make a hot-rolled steel plate,
Tensile strength, characterized in that the cold rolling is made into a cold rolled sheet by subjecting the hot rolled sheet to pickling treatment and then cold rolling at a cold pressure ratio of 10 to 80%. A manufacturing method of high-strength cold-rolled steel sheets with a high proportional limit at 980 MPa or higher.   10. The high-strength cold-cooled sheet according to claim 9, wherein the cold-rolled sheet is further subjected to a continuous annealing treatment having a soaking temperature of 600 to 800 ° C. and a soaking time of 300 s or less to obtain a cold-rolled annealed sheet. A method for producing a thin steel sheet.   During the cooling of the continuous annealing treatment, it is immersed in a galvanizing bath having a bath temperature of 420 to 500 ° C., and is subjected to a hot dip galvanizing treatment to form a hot dip galvanized layer on the steel plate surface to obtain a plated steel plate. The manufacturing method of the high intensity | strength cold-rolled thin steel plate of Claim 10.   The method for producing a high-strength cold-rolled steel sheet according to claim 11, wherein after the hot-dip galvanizing treatment, an alloying treatment of a plating layer is further performed.

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