JP2015147957A - High-strength thin steel sheet excellent in shear-plane delayed-fracture resistance characteristic, and production method thereof - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high-strength thin steel sheet that is excellent in a shear-plane delayed-fracture resistance characteristic.SOLUTION: The thin steel sheet has a composition in which, in mass%, C: 0.100 to 0.250%, Si: 0.3% or less, Mn: 0.1 to 2.0%, Al: 0.10% or less, N: 0.010% or less, and V: 0.400 to 1.00%, are contained, and the C, Mn, N and V are contained by adjusting them such that C-(12/51)×(V-(51/14)×N)≥0.013 and Mn/{V-(51/14)×N}≤2.0 are satisfied, and a structure in which the ferrite phase occupying 95% or more in area ratio is the main phase; the main phase ferrite has a phase in which the ratio of average ferrite grain size in a rolling direction dL and a sheet thickness direction dN, dN/dL, is 0.5 or more, and the average grain size is 5 μm or less, and in which, within the grain, a deposition of less than 10 nm is deposited by 1.0×10pieces/μmor more; and thereby a thin steel sheet having high-strength of the yield ratio of 0.80 or more and the tensile strength TS: 1100 MPa or more, and having an improved shear-plane delayed-fracture resistance characteristic, is obtained.

Description

  The present invention is suitable for structural members used in automobile pillars, side sills, frame members such as members and their reinforcing members, automobile door impact beams, vending machines, desks, home appliances / OA equipment, building materials, etc. The present invention also relates to a high-strength thin steel sheet having excellent delayed fracture resistance 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, the structural member has an advantage that the collision absorption energy can be increased or the plastic deformation can be suppressed by using a high-strength steel plate. However, since the increase in strength makes it easy to cause destruction due to hydrogen penetration, so-called delayed fracture, there is a problem that it becomes difficult to use a high-strength steel sheet for parts used in an environment where hydrogen penetrates. In particular, since delayed fracture occurs from the shear plane, the risk of delayed fracture increases when the end face of the component remains sheared.

  For example, Patent Document 1 describes a high-strength cold-rolled steel sheet and a high-strength hot-dip galvanized steel sheet as steel sheets with improved delayed fracture resistance. The steel sheet described in Patent Document 1 is, by weight, C: 0.05 to 0.3%, Si: 0.3 to 1.6%, Mn: 4.0 to 7.0%, Cr: 0.01 to 0.1%, Ni: 0.02 to 0.1%, Ti : 0.005 to 0.03%, B: 5 to 30ppm, Sb: 0.01 to 0.03%, S: 0.008% or less, Al: 0.5 to 2.0%, Al composition increased, and 40 to 50% tempering High-strength cold-rolled steel sheet that has martensite, 20-40% retained austenite, and a structure composed of the remaining ferrite, and combines high strength of 980 MPa or more, excellent delayed fracture resistance, and excellent ductility. And a high-strength galvanized steel sheet. The steel sheet described in Patent Document 1 can be used not only for bending work but also for general drawing work, and can be expected to reduce the weight of the automobile body and ensure safety.

  Patent Document 2 describes a method for producing a hot-pressed plated steel sheet. In the technique described in Patent Document 2, in mass%, C: 0.01 to 0.5%, Si: 0.05 to 2%, Mn: 0.1 to 3%, Ti: 0.005 to 1%, Nb: 0.01 to 1%, V: 0.01-1%, Mo: 0.01-1%, W: 0.005-1%, Cu: 0.01-3%, Zr: 0.005-0.1%, Y: 0.005-0.5%, Mg: 0.005-1%, La : 0.005 to 0.1%, Ce: 0.005 to 0.1% When annealing and plating a steel sheet with a composition, the amount of non-diffusible hydrogen in the steel sheet before putting it in the annealing furnace is reduced, and annealing is performed. By annealing at 650 ° C or higher and below the specified temperature related to the transformation point in an atmosphere with a hydrogen concentration of 10% or lower and a dew point of 0 ° C or lower in the furnace, and then plating the surface of the steel sheet mainly with aluminum or zinc It is said that a hot-pressed plated steel sheet having excellent delayed fracture resistance can be obtained.

  In Patent Document 3, C: 0.15-0.25%, Si: 1.0-3.0%, Mn: 1.5-2.5%, P: 0.05% or less, S: 0.02% or less, Al: 0.01-0.05%, N: It has a composition containing less than 0.005%, a tempered martensite phase with a volume ratio of 40 to 85%, and a structure containing a ferrite phase with a volume ratio of 15 to 60%, and a tensile strength of 1320 MPa or more. Cold rolled steel sheet is described. The technology described in Patent Document 3 contains a high amount of C and Si, increases the hardness of tempered martensite and ferrite, ensures a high strength of 1320 MPa or higher, and contains almost no dislocations. By precipitating the ferrite phase, the dislocation density in the metal structure is greatly reduced compared to the martensite single phase structure, the hydrogen trap sites are reduced, and the amount of hydrogen penetrating into the steel is greatly reduced. It is going to improve delayed fracture resistance.

Special Table 2011-523442 JP 2012-41597 JP 2012-12642

  However, the steel sheet manufactured by the technique described in Patent Document 1 has a problem that the corrosion resistance deteriorates because it contains a large amount of Mn. Moreover, the technique described in Patent Document 2 is a plated steel sheet for members manufactured by hot pressing, and is based on the premise that a hot press heated to an austenite region is used. There is no mention of delayed fracture resistance of surfaces. Further, the steel sheet manufactured by the technique described in the cited document 3 can reduce the hydrogen trap sites, but the structure is a composite structure of tempered martensite and ferrite, and the ferrite phase and martensite are sheared during shearing. Cracks occurred at the interface of the site phase, and therefore there was a problem that improvement in delayed fracture resistance of the shear plane could not be expected.

  The present invention advantageously solves the problems of the prior art, yield ratio (= yield strength YP / tensile strength TS) is 0.80 or more, tensile strength TS: 1100MPa or more, and delayed fracture resistance of the shear plane. An object is to provide an excellent high-strength thin steel sheet and a method for producing the same. The term “high strength” as used herein refers to the case where the tensile strength TS is 1100 MPa or more, preferably 1180 MPa or more, more preferably 1320 MPa or more, and the “thin steel plate” preferably has a thickness of 2.0. It shall mean the case where it is less than mm.

  In order to achieve the above-mentioned object, the present inventors diligently studied various factors that affect the delayed fracture resistance of the sheared surface. As a result, the inventors came up with the idea of making the structure a ferrite single phase. This is because if a soft phase such as a ferrite phase and a hard phase such as a martensite phase are present in combination, microcracks are generated at the interface during shearing, inducing hydrogen cracking, and the delayed fracture resistance of the sheared surface is reduced. To do.

  Furthermore, by further study by the present inventors, ferrite is made into equiaxed and fine crystal grains, and by depositing a large number of fine precipitates in the ferrite grains, both delayed fracture resistance and high strength are achieved. It has been found that it can be remarkably improved. This is because coarse and stretched grains concentrate processing strain in the vicinity of the grain boundary during shearing, becoming a hydrogen adsorption source, and tend to cause embrittlement due to hydrogen. It was found that if the grains are equiaxed and fine, concentration of processing strain near the grain boundary can be avoided, and embrittlement caused by hydrogen can be suppressed.

Further, the fine precipitates contribute to high strength and also effectively contribute as hydrogen trap sites. The inventors have found that by precipitating a large number of fine precipitates, high strength can be secured, hydrogen adsorption to dislocations introduced during shearing can be suppressed, and deterioration in delayed fracture resistance can also be suppressed.
Next, the results of experiments conducted by the present inventors and serving as the basis of the present invention will be described.

  A 30 × 100 mm plate was cut out by shearing from a hot-rolled steel plate with various changes in the manufacturing conditions and the crystal grain size of the ferrite single-phase structure varied, and then U-bent with a punch with a bending radius of 5 mm. As shown in FIG. 2, a sample in which the springback was tightened with a bolt was prepared, a delayed fracture test was performed, and the delayed fracture resistance was evaluated. The delayed fracture test was a test in which the sample shown in FIG. 2 was immersed in a 0.1N hydrochloric acid buffer (liquid temperature: room temperature) and the time until cracking (breakage time) was measured. The immersion time was up to 200h, and the samples that did not crack by 200h were not cracked.

  The obtained results are shown in FIG. 1 in relation to the destruction time and dN / dL. DN is the average ferrite grain size in the plate thickness direction, and dL is the average ferrite grain size in the rolling direction. From FIG. 1, it can be seen that as dN / dL approaches 1, that is, as the ferrite grains approach the equiaxes, the time to failure is extended. From the examination conducted separately, it was concluded that the delayed fracture resistance of the sheared surface was good if the fracture time was 80h or more in the above immersion test. It was found that dN / dL must be 0.5 or more in order to improve.

  Furthermore, the present inventors examined various factors in the production of steel sheets that affect ferrite crystal grains. As a result, in order to obtain equiaxed and fine ferrite grains, rough rolling, especially in hot rolling, should be performed under conditions that terminate the rolling at a high pressure reduction rate of 82% or more and a high temperature of 950 ° C or more. In addition, it was found that it is important to perform finish rolling under the condition that the rolling is finished at a high pressure ratio of 82% or more and a high temperature of 850 ° C. or more.

The present invention has been completed based on such findings and further studies. That is, the gist of the present invention is as follows.
(1) By mass%, C: 0.100 to 0.250%, Si: 0.3% 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 , V: 0.400 to 1.00% are included, and C, Mn, N, and V are represented by the following formulas (1) and (2):
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
(Here, C, V, N, Mn: Content of each element (mass%))
The composition comprising the balance Fe and inevitable impurities and a ferrite phase with an area ratio of 95% or more as a main phase, the ferrite phase having an average ferrite grain diameter dN in the thickness direction The ratio of dN / dL to the average ferrite grain size dL in the rolling direction is 0.5 or more, the average grain size defined by (2 x dL x dN) / (dL + dN) is 5 μm or less, and a precipitate of less than 10 nm It has a microstructure with a precipitation density of 1.0 × 10 ⁵ pieces / μm ³ or more, a tensile strength TS: 1100 MPa or more, and a yield ratio YR: 0.8 or more. Excellent high-strength thin steel sheet.
(2) In (1), in addition to the above composition, in terms of mass%, Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005 ˜0.600% of one or more selected from the following (3) to C, Mn, N, V depending on the relationship of Ti, N, S instead of the above formula (1) (5) Formula
When Ti− (48/14) × N− (48/32) × S ≧ 0,
C− (12/51) × V− (12/48) × (Ti− (48/14) × N− (48/32) × S) − (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (3)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N ≧ 0,
C− (12/51) × V− (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (4)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N <0,
C− (12/51) × (V− (51/14) × (N− (14/48) × Ti)) − (12/93) × Nb− (12/96) × Mo− (12/181 ) × Ta− (12/184) × W ≧ 0.013 (5)
(Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: content of each element (mass%))
Any one of the following, in place of the above formula (2), the following formula (6) or (7) according to the relationship between Ti and N
When Ti− (48/14) × N ≧ 0, Mn / V ≦ 2.0 (6)
When Ti− (48/14) × N <0,
Mn / (V− (51/14) × (N− (14/48) × Ti)) ≦ 2.0 (7)
(Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: content of each element (mass%))
The high-strength thin steel sheet according to claim 1, wherein the high-strength steel sheet is adjusted so as to satisfy the requirements.
(3) A high-strength 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, in addition to mass, one kind of Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cu: 0.01 to 1.0%, or A high-strength thin steel sheet containing two or more types.
(5) The high strength 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), in addition to the above composition, the composition further contains one or two of Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01% by mass%. High strength thin steel sheet.
(7) A high-strength thin steel sheet excellent in delayed fracture resistance of a shear surface, comprising a plating layer on the surface of the high-strength thin steel sheet according to any one of (1) to (6).
(8) The steel material is subjected to hot rolling consisting of rough rolling and finish rolling, and then cooled, wound, and made into a high-strength steel sheet. 0.250%, Si: 0.3% or less, Mn: 0.1-2.0%, P: 0.05% or less, S: 0.030% or less, Al: 0.10% or less, N: 0.010% or less, V: 0.400-1.00%, and , C, Mn, N, and V are expressed by the following equations (1) and (2)
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
(Here, C, V, N, Mn: Content of each element (mass%))
The steel material having a composition comprising the balance Fe and unavoidable impurities is adjusted so as to satisfy the requirements, and the rough rolling is performed at a rough rolling total rolling reduction of 82% or more and a rough rolling finish temperature: 950 ° C. or more. The finish rolling is a rolling in which the finish rolling total reduction ratio is 92% or more and the finish rolling finish temperature is 850 ° C. or more. After the finish rolling, cooling in the temperature range from the finish finish finish to 700 ° C. A method for producing a high-strength thin steel sheet excellent in delayed fracture resistance of a sheared surface, characterized by cooling at an average speed of 30 ° C / s or more and performing the winding at a winding temperature of 500 ° C or more.
(9) In (8), in addition to the composition of the steel material, further, in mass%, Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600% , W: One or two or more selected from 0.005 to 0.600%, C, Mn, N, and V are changed according to the relationship of Ti, N, and S, instead of the formula (1). Equations (3) to (5)
When Ti− (48/14) × N− (48/32) × S ≧ 0,
C− (12/51) × V− (12/48) × (Ti− (48/14) × N− (48/32) × S) − (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (3)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N ≧ 0,
C− (12/51) × V− (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (4)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N <0,
C− (12/51) × (V− (51/14) × (N− (14/48) × Ti)) − (12/93) × Nb− (12/96) × Mo− (12/181 ) × Ta− (12/184) × W ≧ 0.013 (5)
(Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: content of each element (mass%))
In addition to the above equation (2), the following (6) or the following (7) equation depending on the relationship between Ti and N:
When Ti− (48/14) × N ≧ 0, Mn / V ≦ 2.0 (6)
When Ti− (48/14) × N <0,
Mn / (V− (51/14) × (N− (14/48) × Ti)) ≦ 2.0 (7)
(Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: content of each element (mass%))
The method for producing a high-strength thin steel sheet according to claim 8, wherein the content is adjusted so as to satisfy the requirements.
(10) The method for producing a high-strength thin steel sheet according to (8) or (9), further comprising B: 0.0003 to 0.0050% by mass% in addition to the composition of the steel material.
(11) In any one of (8) to (10), in addition to the composition of the steel material, further in mass%, Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cu: 0.01 to 1.0% A method for producing a high-strength thin steel sheet, comprising one or more selected from among the above.
(12) In any one of (8) to (11), in addition to the composition of the steel material, in addition to the mass, Sb: 0.005 to 0.050% is contained. Method.
(13) In any one of (8) to (12), in addition to the composition of the steel material, 1% selected from Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01% by mass% A method for producing a high-strength thin steel sheet, comprising seeds or two kinds.
(14) In any one of (8) to (13), when the hot-rolled thin steel sheet after the winding is subjected to pickling and annealing, the annealing has a soaking temperature in a range of 750 ° C. or less. The temperature range from 500 ° C. to the soaking temperature is heated at an average heating rate of 1 ° C./s or more, soaking is carried out at a soaking time at the soaking temperature of 300 s or less, and after soaking, 1 High strength, characterized in that it is cooled to 500 ° C. at a cooling rate of at least ° C./s, and is subjected to a plating treatment immersed in a galvanizing bath at a bath temperature of 420 to 500 ° C. during the cooling process of the annealing. Manufacturing method of thin steel sheet.
(15) In (14), after the plating process, the heating temperature is further reheated to 460 to 600 ° C., and the alloying process of the plating layer which is held at the heating temperature for 1 s or more is performed. A method of manufacturing a steel sheet.
(16) The method for producing a high-strength thin steel sheet according to any one of (8) to (13), further comprising providing a processing of a sheet thickness reduction rate of 0.1 to 3.0% after the winding.
(17) The method for producing a high-strength thin steel sheet according to (14), wherein after the plating treatment, further processing with a thickness reduction ratio of 0.1 to 3.0% is applied.
(18) The method for producing a high-strength thin steel sheet according to (15), wherein after the alloying treatment, further processing with a sheet thickness reduction rate of 0.1 to 3.0% is applied.

  According to the present invention, a high-strength sheet steel having a yield ratio of 0.80 or more, a tensile strength TS of 1100 MPa or more, and excellent delayed fracture resistance on a shear surface can be easily and stably produced. Yes, and it has a remarkable industrial effect.

5 is a graph showing the relationship between the fracture time during immersion in 0.1N hydrochloric acid buffer and the ratio dN / dL of the ferrite grain thickness direction average particle diameter dN and rolling direction average particle diameter dL. It is explanatory drawing which shows typically the U bending test piece (immersion test piece) used by the delayed fracture test.

The high-strength thin steel sheet of the present invention is a steel sheet having a high strength of yield ratio YR: 0.80 or more and tensile strength TS: 1100 MPa or more, in mass%, C: 0.100 to 0.250%, Si: 0.3% or less, Mn : 0.1-2.0%, P: 0.05% or less, S: 0.030% or less, Al: 0.10% or less, N: 0.010% or less, V: 0.400-1.00%, and C, Mn, N, V, Next (1) and next (2)
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
(Here, C, V, N, Mn: Content of each element (mass%))
In order to satisfy the above, it is contained and has a composition consisting of the balance Fe and inevitable impurities.

First, the reasons for limiting the composition will be described. Hereinafter, mass% in the composition is simply expressed as% unless otherwise specified.
C: 0.100 to 0.250%
C is an element that combines with carbide-forming elements such as 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. In order to obtain such an effect, the content of 0.100% or more is required. In addition, Preferably it is 0.120% or more, More preferably, it is 0.150% or more. On the other hand, a large content exceeding 0.250% suppresses ferrite transformation and promotes transformation to bainite and martensite. For this reason, fine carbide formation with V is suppressed. Moreover, a large amount of C content also deteriorates weldability. For this reason, C was limited to 0.250% or less. In addition, Preferably it is 0.200% or less. Therefore, C is limited to the range of 0.100 to 0.250%.

Si: 0.3% or less
Si is a ferrite-forming element, and if contained in a large amount, Si has an adverse effect of promoting ferrite transformation in cooling after hot rolling, raising 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 unfavorable part such as unplated part is adversely affected. In order to avoid such adverse effects, it is preferable to reduce Si as much as possible in the present invention, but such adverse effects are acceptable up to 0.3%. For this reason, Si was limited to 0.3% 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. The impurity level of Si is about 0.005%.

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 V is suppressed. 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. Therefore, in the present invention, it is desirable to reduce P as much as possible, 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 deoxidizing agent. 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 0.05% or less.

N: 0.010% or less
N forms coarse nitrides at high temperatures with V, and contributes less to strength, while reducing the contribution of V to higher strength. 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.

V: 0.400 to 1.00%
V combines with C to form fine carbides, contributing to high strength. In order to obtain such an effect, a content of 0.0400% or more is required. On the other hand, even if the content exceeds 1.00%, an effect commensurate with the content cannot be expected, which is economically disadvantageous. For this reason, V was made into 1.00% or less. In addition, Preferably it is 0.800% or less, More preferably, it is 0.600% or less. For this reason, V is limited to the range of 0.400 to 1.00%.

In the present invention, C, Mn, N, and V are within the above range, and the following formula (1) and the following formula (2)
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
(Here, C, V, N, Mn: Content of each element (mass%))
The content is adjusted so as to satisfy.

C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Equation (1) defines the relationship among the C, V, and N contents for precipitating fine VC and achieving the desired high strength. When the amount of V excluding V that forms nitrides by combining with N (effective V amount) increases with respect to the amount of C in terms of atomic ratio, the amount of solute V increases, which contributes to higher strength of VC. The amount of precipitation is small and the desired high strength cannot be achieved. For this reason, {C− (12/51) × (V− (51/14) × N)} is limited to 0.013 or more. If {C− (12/51) × (V− (51/14) × N)} is less than 0.013, the desired increase in strength cannot be achieved. In addition, Preferably it is 0.020 or more. On the other hand, if {C− (12/51) × (V− (51/14) × N)} exceeds 0.050, a large amount of surplus C is generated, forming cementite, and the workability is significantly reduced. Let For this reason, {C− (12/51) × (V− (51/14) × N)} is preferably 0.050 or less. For this reason, C, V, and N are adjusted to satisfy the expression (1).

Mn / {V− (51/14) × N} ≦ 2.0 (2)
Equation (2) defines the relationship among the contents of Mn, V, and N in order to finely precipitate V carbide (VC) in ferrite and achieve a desired high strength. If the amount of Mn is too large compared to the amount of V excluding V that combines with N to form nitride (effective amount of V), the ferrite transformation start temperature becomes too low, and the runout table during hot rolling After being wound up or coiled, V carbide precipitates in austenite, not in ferrite, and becomes coarse carbide, making it impossible to achieve the desired high strength. Such a phenomenon can be prevented by setting Mn / {V− (51/14) × N} to 2.0 or less. For this reason, in the present invention, Mn / {V− (51/14) × N} is limited to 2.0 or less. In addition, Preferably it is 1.5 or less, More preferably, it is 1.2 or less, More preferably, it is 1.0 or less. The lower limit is not particularly limited, but is preferably 0.1 or more from the viewpoint of stabilizing the strength.

  The above-described components are basic components. In the present invention, Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005 1 or 2 or more types selected from ˜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 1.0 1 or 2 or more selected from%, 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% Or 2 types can be contained.

One or more selected from Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005-0.600%
Ti, Nb, Mo, Ta, and W are elements that contribute to high strength by precipitation strengthening due to fine carbide precipitation, and can be selected as needed and contained in one or more.
Ti, Nb, Mo, Ta, and W all form fine carbides and contribute to high strength. Ti forms a composite carbide together with V and contributes to high strength. In order to obtain such an effect, it is necessary to contain Ti: 0.005% or more, Nb: 0.005% or more, Mo: 0.005% or more, Ta: 0.005% or more, W: 0.005% or more. Preferably, Ti is 0.050% or more, 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 Ti: 0.600%, Nb: 0.600%, Mo: 0.600%, Ta: 0.600%, W: 0.600%, even if contained in a large amount, the effect is saturated and an effect commensurate with the content cannot be expected, Economic disadvantage. Moreover, when a large amount is contained, coarse precipitates are formed, and there is an adverse effect of deteriorating the steel sheet characteristics. More preferably, Ti is 0.300% or less, 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, in the case of inclusion, Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005-0.600% It is preferable to limit.

When one or more selected from Ti, Nb, Mo, Ta, and W are contained, the formula that C, N, V, etc. should be satisfied instead of the formula (1) described above is , (3) to (5).
Since Ti forms nitrides and further sulfides in preference to V, the amount of {Ti− (48/14) × N− (48/32) × S} and {Ti− (48/14) Depending on the amount of × N}, the formula to be applied is changed, and C, N, and V are adjusted so as to satisfy the formula to achieve a desired high strength.

When Ti− (48/14) × N− (48/32) × S ≧ 0, Ti bonded to C remains, so the following equation (3) is applied.
C− (12/51) × V− (12/48) × (Ti− (48/14) × N− (48/32) × S) − (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (3)
Apply.

In addition, when Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N ≧ 0, Ti bonding to C is deficient. 4) Formula
C− (12/51) × V− (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (4)
Apply.

In addition, when Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N <0, Ti bonding to C is deficient, but N is Since it remains, the following equation (5)
C− (12/51) × (V− (51/14) × (N− (14/48) × Ti)) − (12/93) × Nb− (12/96) × Mo− (12/181 ) × Ta− (12/184) × W ≧ 0.013 (5)
Apply.

In addition, when one or more selected from Ti, Nb, Mo, Ta, W are contained, Mn is substituted for the above formula (2) according to the relationship between Ti and N. , Change the formula that V should satisfy.
Since Ti forms nitrides preferentially over V, the amount of N combined with V changes according to the amount of Ti, and the amount of V excluding V combined with N changes. Therefore, according to the amount of {Ti− (48/14) × N}, the formula to be applied is changed instead of the formula (2), and Mn and V are adjusted so as to satisfy the formula. Achieve strength.

When Ti− (48/14) × N ≧ 0, N that binds to V is deficient, so instead of equation (2), the following equation (6)
Mn / V ≦ 2.0 (6)
Apply. In addition, when Ti− (48/14) × N <0, N that binds to V remains, so instead of the formula (2), the following formula (7)
Mn / (V− (51/14) × (N− (14/48) × Ti)) ≦ 2.0 (7)
Apply.

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 the improvement of secondary work brittleness resistance. In order to acquire such an effect, it is preferable to contain 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, it is preferable to limit 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 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 is an element that segregates on the surface of the slab during hot rolling, prevents the slab from 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 it is contained in a large amount exceeding 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 those described above is Fe and inevitable impurities. Examples of impurities include Sn, Mg, Co, As, Pb, Zn, O, and the like, but these impurity elements can be allowed without any problem in characteristics if the total is 0.5% or less.
Further, the high-strength thin steel sheet of the present invention has the above-described composition, and further comprises a ferrite phase having an area ratio of 95% or more as a main phase and the main phase and a second phase having an area ratio of 0 to 5%. Have an organization.

Main phase: 95% or more ferrite phase in area ratio In the present invention, the main phase is a ferrite phase in order to improve moldability. The “main phase” here is a phase that occupies 95% or more in area ratio. 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 the transformation, and the yield strength decreases. In addition, cracks are generated at the interface between the ferrite and the low-temperature transformation phase during shearing, and the delayed fracture resistance of the sheared surface is significantly reduced. For this reason, the ferrite phase is 95% or more in area ratio. The area ratio is preferably 98% or more, and more preferably 100%. From the viewpoint of delayed fracture resistance of the shear plane, the second phase other than the ferrite phase is preferably 0% in area ratio. Even if it exists, the area ratio is 5% or less.

The ferrite phase, which is the main phase of the high strength thin steel sheet of the present invention, has a ratio dN / dL between the average ferrite grain size dN in the plate thickness direction and the average ferrite grain size dL in the rolling direction of 0.5 or more, and an average grain size (2 × dL × dN) / (dL + dN) is 5 μm or less, and among the precipitates precipitated in the ferrite phase, the number density of precipitates of less than 10 nm is 1.0 × 10 ⁵ pieces / μm ³ or more.
Ratio of average ferrite grain diameter dN in the plate thickness direction and average ferrite grain diameter dL in the rolling direction dN / dL: 0.5 or more In the expanded grain, processing strain concentrates near the grain boundary during shearing, and becomes a hydrogen adsorption source , Embrittlement caused by hydrogen occurs. As the stretched grain approaches the equiaxed grain, concentration of processing strain near the grain boundary can be avoided, and embrittlement due to hydrogen can be suppressed. As shown in FIG. The fracture time that occurs is longer and the delayed fracture resistance is improved. If dN / dL is 0.5 or more, desired delayed fracture resistance can be secured. For this reason, dN / dL was limited to 0.5 or more. In addition, Preferably it is 0.6 or more, More preferably, it is 0.7 or more, More preferably, it is 0.8 or more. Here, “dN” and “dL” are the average ferrite grain size dN in the plate thickness direction and the average ferrite grain size dL in the rolling direction, and the cross section in the rolling direction is polished and corroded to obtain a scanning electron microscope (magnification: 1000 times), the grain size in the rolling direction (cutting length) and the grain size in the thickness direction (cutting length) are obtained by the cutting method, and the obtained values are obtained by arithmetic averaging. Value.

Average grain size of ferrite (2 × dL × dN) / (dL + dN): 5 μm or less In coarsened grains, processing strain concentrates near the grain boundary during shearing, becoming a hydrogen adsorption source, and embrittlement caused by hydrogen occurs. Arise. Therefore, in order to prevent embrittlement due to hydrogen and suppress the occurrence of delayed fracture, it is necessary to use fine crystal grains. According to the knowledge of the present inventors, when the average particle diameter of ferrite is 5 μm or less, embrittlement due to hydrogen can be prevented and a decrease in delayed fracture resistance can be suppressed. For this reason, the average particle diameter of ferrite was limited to 5 μm or less. In addition, Preferably it is 4.0 micrometers or less, More preferably, it is 3.0 micrometers or less, More preferably, it is 2.0 micrometers or less. The “average particle size” here is a value calculated by (2 × dL × dN) / (dL + dN) using the average ferrite particle size dN in the plate thickness direction and the average ferrite particle size dL in the rolling direction. Shall be used.

Precipitation density of precipitates less than 10 nm in the ferrite phase: 1.0 × 10 ⁵ pieces / μm ³ or more In order to achieve high strength by precipitation strengthening, it is necessary to disperse a large amount of fine precipitates. In the present invention, in order to secure a desired high strength of tensile strength TS: 1100 MPa or more, fine precipitates of less than 10 nm are 1.0 × 10 ⁵ pieces / μm ³ or more, preferably 2.0 × 10 ⁵ pieces / μm ³ or more, more preferably 3.0 × 10 ⁵ particles / μm ³ or more, and further preferably 4.0 × 10 ⁵ particles / μm ³ or more is dispersed in a large amount. The coarse precipitate hardly contributes to the increase in strength and also becomes a starting point of hydrogen embrittlement.

Further, the particle diameter of the fine precipitate is a maximum diameter of less than 10 nm, preferably less than 5 nm, and more preferably less than 3 nm.
In the high strength 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, a preferred method for producing the high strength thin steel sheet of the present invention will be described.
In the present invention, the steel material having the composition described above is subjected to hot rolling consisting of rough rolling and finish rolling, and then cooled, wound, and made into a hot rolled thin steel plate.
The manufacturing method of the steel material is not particularly limited, but the molten steel having the above composition is melted by a conventional melting method such as a converter, and a slab such as a slab by a conventional casting method such as continuous casting. It is preferable to use a steel material.

In the case where 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.
In addition, it is preferable that heating temperature shall be 1200 degreeC or more and less than a 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. Below 1200 ° C, the carbide-forming elements cannot be completely dissolved. The heating of the steel material is preferably maintained at 1200 ° C. or higher for 10 minutes or longer. More preferably at 1250 ° C. or more, 10 min or more, more preferably at 1250 ° C., 30 min or more, more preferably at 1300 ° C. or more, 10 min or more, more preferably at 1350 ° C. or more, 10 min or more, more preferably at 1400 ° C. or more, 10 min or more. That's fine.

The heated steel material is then subjected to hot rolling consisting of rough rolling and finish rolling to form a hot rolled thin steel plate. Unless otherwise specified, the temperature is the steel sheet surface temperature.
In the present invention, in order to form fine and equiaxed grained ferrite grains, rough rolling is adjusted to an appropriate condition. Rough rolling is rolling with a total rolling reduction of 82% or more and a rough rolling end temperature of 950 ° C. or more.

Rough rolling total rolling reduction: 82% or more In order to obtain fine and equiaxed grains, the total rolling reduction in rough rolling is limited to 82% or more in the present invention. If the total rolling reduction ratio is less than 82%, the amount of reduction is too small to obtain a desired thin steel sheet having fine and equiaxed ferrite crystal grains. It is preferably 84% or more, more preferably 86% or more, and further preferably 88% or more. The upper limit of the total rolling reduction is not particularly limited, but is 95% or less from the viewpoint of operability.

Coarse rolling end temperature: 950 ° C. or higher In order to equiax the ferrite grains, it is necessary to maintain the rough rolling end temperature as high as 950 ° C. or higher. When the end temperature of the rough rolling is lowered, it tends to be expanded grains, so the end temperature of the rough rolling is limited to 950 ° C. or higher. In addition, Preferably it is 980 degreeC or more.
In the present invention, after the above-described rough rolling is performed, the finish rolling is a rolling with a finish rolling total reduction of 92% or more and a finish rolling finish temperature of 850 ° C. or more.

Finish rolling total reduction ratio: 92% or more In order to obtain fine and equiaxed grains, the total rolling reduction in finish rolling is limited to 92% or more in the present invention. If the total reduction amount is less than 92%, the reduction amount is too small to obtain a desired thin steel plate having fine and equiaxed ferrite crystal grains. In addition, Preferably it is 94% or more, More preferably, it is 96% or more. The upper limit of the total rolling reduction is not particularly limited, but is 98% or less from the viewpoint of operability.

Finish rolling end temperature: 850 ° C. or higher If the finish rolling end temperature is low, ferrite transformation is promoted by cooling after rolling, precipitation of carbide is promoted, and coarse carbide is formed, so that desired precipitation strengthening cannot be expected. Further, when the finish rolling finish temperature is low enough to be in the ferrite region, coarse carbides precipitate due to strain-induced precipitation, and ferrite grains become non-axial. 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 and making a hot-rolled thin steel sheet, it is further cooled and wound. Cooling is a process of cooling at 30 ° C./s or higher at an average cooling rate from the end of finish rolling to 700 ° C. The coiling temperature is 500 ° C or higher.
Average cooling rate from finish rolling to 700 ° C: 30 ° C / s or more When cooling from finish rolling to 700 ° C is slow at an average of less than 30 ° C / s, ferrite transformation is promoted and carbide precipitates greatly. To do. 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 suppression.

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 wound hot-rolled steel sheet may be pickled, annealed, and further plated.
Annealing is a pre-treatment for plating the steel sheet, soaking temperature is 750 ° C or lower, preferably 600 to 750 ° C, and an average temperature range from 500 ° C to soaking temperature is 1 ° C. Heating at a heating rate of not less than / s, soaking at the soaking temperature for a holding time of not more than 300 s, and cooling to 500 ° C. at a cooling rate of not less than 1 ° C./s is preferable.

Soaking temperature: 750 ° C. or less When the soaking temperature exceeds 750 ° C. and becomes high, fine precipitates are coarsened and crystal grains are also coarsened. For this reason, the soaking temperature during annealing is limited to 750 ° C. or lower, preferably 600 ° C. or higher and 750 ° C. or lower. In addition, Preferably it is 720 degrees C or less, More preferably, it is 700 degrees C or less. In addition, the lower limit of the soaking temperature should be able to ensure a temperature equal to or higher than the zinc plating bath temperature (420 to 500 ° C), but if the surface property of the plating needs to be good, the soaking temperature is 600 ° C or more. It is preferable that More preferably, it is 650 ° C. or higher.

Heating rate during annealing: 1 ° C / s or more on average from 500 ° C to soaking temperature When applying hot dip galvanizing, if the heating rate from 500 ° C to soaking temperature is slow at less than 1 ° C / s, Fine precipitates are coarsened and the desired properties cannot be secured. Therefore, the average heating rate from 500 ° C to the soaking temperature was limited to 1 ° C / s or more. In addition, Preferably it is 5 degrees C / s or more, More preferably, it is 10 degrees C / s or more. If the heating rate exceeds 1000 ° C./s and becomes too high, it will be difficult to control the soaking temperature. The average is more preferably 300 ° C./s or less, further preferably 100 ° C./s or less, and still more preferably 50 ° C./s or less.

Soaking time: 300 s or less If the soaking time during annealing exceeds 300 s and becomes a long time, fine precipitates become coarse. For this reason, the soaking time is limited to 300 s or less, more preferably 200 s or less, and still more preferably 100 s or less. If the soaking time is less than 1 s, a desired soaking state cannot be secured.
Average cooling rate from soaking temperature to 500 ° C (plating bath temperature): 1 ° C / s or more When cooling in the temperature range from soaking temperature to 500 ° C (plating bath temperature) is slow, fine precipitates become coarse The desired characteristics cannot be ensured. For this reason, the average cooling rate from the soaking temperature to the plating bath was limited to 1 ° C./s or more. The temperature is preferably 3 ° C./s or more, more preferably 5 ° C./s or more, and further preferably 10 ° C./s or more. From the viewpoint of preventing the coarsening of the precipitate, an average cooling rate of 100 ° C./s or less is sufficient.

The plating process is a process of immersing in a zinc plating bath having a bath temperature of 420 to 500 ° C. in the course of annealing cooling.
Further, in the present invention, after the plating treatment, as the alloying treatment of Zn and Fe, it is heated to a temperature of 460 to 600 ° C., and is kept for 1 s or more, preferably 10 s or less from the viewpoint of preventing coarsening of precipitates. (Alloying treatment) may be applied. In addition, since alloying advances too much and the plating becomes brittle when the reheating temperature is increased, the temperature is preferably in the range of 460 to 600 ° C., more preferably 570 ° C. or less.

  Light processing may be applied to the thin steel sheet after hot rolling, after annealing, after plating, or after alloying. By adding light processing to a thin steel plate, it is possible to increase movable dislocations and improve formability and shape freezeability. As the light processing to be applied, it is preferable that the thickness reduction ratio is 0.1 to 3.0%. If the plate thickness reduction rate is less than 0.1%, the light processing effect described above cannot be obtained. For this reason, it is preferable to limit the sheet thickness reduction rate in light processing to 0.1% or more. In addition, More preferably, it is 0.3% or more. On the other hand, when the plate thickness reduction rate exceeds 3.0%, dislocations are difficult to move due to dislocation interaction, and formability is deteriorated. For this reason, it is preferable that the plate | board thickness reduction | decrease rate in light processing shall be 3.0% or less. In addition, More preferably, it is 2.0% or less, More preferably, it is 1.0% or less. For this reason, when light processing is performed, it is preferable to limit the processing to a plate thickness reduction rate of 0.1 to 3.0%. Here, the light processing means is preferably processing by a rolling roll, processing by tension, or combined processing combining rolling and tension processing.

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 under the conditions shown in Table 2 to obtain hot-rolled steel sheets having the thicknesses shown in Table 2. Some hot-rolled steel sheets were further annealed under the conditions shown in Table 2, or further plated, or further alloyed with a plated layer, or further light worked.
A test piece was collected from the obtained steel sheet, and subjected to structure observation, tensile test, and delayed fracture test, and evaluated for strength and delayed fracture resistance. 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). The structure of the 300 × 300 μm ² region was imaged, and the area ratio of ferrite was calculated. In addition, the region of 100 × 100 μm ² was observed with a scanning electron microscope (magnification: 1000 times), imaged, and the grain size of the ferrite obtained was measured by measuring the ferrite grain size in the rolling direction and the plate thickness direction by a cutting method. The diameters were arithmetically averaged to obtain a ferrite average particle diameter dL in the rolling direction and a ferrite average particle diameter dN in the plate thickness direction.

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.
(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.
(3) Delayed fracture test From the obtained steel plate, a test piece (t x 30 x 100 mm) was sampled by shearing, U-bent with a punch with a bending radius of 5 mm, and the spring back was tightened with bolts as shown in Fig. 2 And a sample for immersion. The obtained sample for immersion was immersed in a 0.1N hydrochloric acid buffer (liquid temperature: room temperature), and the crack occurrence time (destruction time: h) was measured to evaluate delayed fracture resistance. In addition, when no cracking occurred at the immersion time of 200 hours, “no cracking” was set. In the present invention, when the fracture time is 80 h or more due to the above-described delayed fracture, the delayed fracture resistance is excellent.

  The obtained results are shown in Table 3.

  In all of the inventive examples, the yield ratio YR is 0.80 or more, the tensile strength TS is high strength of 1100 MPa or more, the crack occurrence time (fracture time) is 80 h or more in the delayed fracture test, and delayed fracture resistance is achieved. It is an excellent high-strength thin steel sheet. On the other hand, the comparative example out of the scope of the present invention is that the yield ratio YR is less than 0.80 or the tensile strength TS is less than 1100 MPa, and the desired high strength cannot be secured, or the delay test is broken. When the time is less than 80h, delayed fracture resistance is reduced.

Claims (18)

% By mass
C: 0.100 to 0.250%, Si: 0.3% 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, V: 0.400 to 1.00%
And containing C, Mn, N, and V so as to satisfy the following formula (1) and the following formula (2), and the balance Fe and inevitable impurities:
A ferrite phase having an area ratio of 95% or more is a main phase, and the ferrite phase has a ratio dN / dL of an average ferrite particle diameter dN in the sheet thickness direction and an average ferrite particle diameter dL in the rolling direction of 0.5 or more, (2 × dL × dN) / (dL + dN) as defined by an average particle size of 5 μm or less, and a precipitate density of less than 10 nm is 1.0 × 10 ⁵ particles / μm ³ or more, and tensile A high-strength steel sheet with excellent delayed fracture resistance on sheared surfaces, characterized by strength TS: 1100 MPa or more and yield ratio YR: 0.8 or more.
Record
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
Where C, V, N, Mn: Content of each element (% by mass) In addition to the above composition, it is further selected from Ti: 0.005 to 0.600%, Nb: 0.005 to 0.600%, Mo: 0.005 to 0.600%, Ta: 0.005 to 0.600%, W: 0.005 to 0.600%. In addition, one or two or more of C, Mn, N, and V are replaced by the above formula (1), and any one of the following formulas (3) to (5) is selected according to the relationship between Ti, N, and S. Furthermore, it replaces with said (2) Formula, and it adjusts and contains so that the following (6) or (7) Formula may be satisfied according to the relationship of Ti and N, The high of Claim 1 characterized by the above-mentioned. Strength thin steel sheet.
Record
When Ti− (48/14) × N− (48/32) × S ≧ 0,
C− (12/51) × V− (12/48) × (Ti− (48/14) × N− (48/32) × S) − (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (3)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N ≧ 0,
C− (12/51) × V− (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (4)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N <0,
C− (12/51) × (V− (51/14) × (N− (14/48) × Ti)) − (12/93) × Nb− (12/96) × Mo− (12/181 ) × Ta− (12/184) × W ≧ 0.013 (5)
When Ti− (48/14) × N ≧ 0, Mn / V ≦ 2.0 (6)
When Ti− (48/14) × N <0,
Mn / (V− (51/14) × (N− (14/48) × Ti)) ≦ 2.0 (7)
Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: Content of each element (mass%)   The high-strength thin 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, the composition further contains one or more of Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, Cu: 0.01 to 1.0% by mass%. 4. The high-strength thin steel sheet according to any one of 3.   The high strength thin steel sheet according to any one of claims 1 to 4, further comprising Sb: 0.005 to 0.050% by mass% in addition to the composition.   The high content according to any one of claims 1 to 5, further comprising one or two of Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01% by mass% in addition to the composition. Strength thin steel plate.   A high-strength steel sheet excellent in delayed fracture resistance of a shear surface, comprising a plating layer on the surface of the high-strength thin steel sheet according to claim 1. The steel material is subjected to hot rolling consisting of rough rolling and finish rolling, and then cooled, wound, and made into a high-strength thin steel plate.
The steel material,
% By mass
C: 0.100 to 0.250%, Si: 0.3% 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, V: 0.400 to 1.00%
And containing C, Mn, N and V so as to satisfy the following formula (1) and the following formula (2), and having a composition comprising the balance Fe and inevitable impurities ,
The rough rolling is a rolling with a rough rolling total rolling reduction of 82% or more and a rough rolling end temperature: 950 ° C. or more,
The finish rolling is a rolling with a finish rolling total reduction ratio of 92% or more and a finish rolling finish temperature: 850 ° C. or more,
The cooling after the finish rolling is cooled at an average cooling rate of 30 ° C./s or higher in the temperature range from the finish rolling to 700 ° C.
Winding temperature: The manufacturing method of the high strength thin steel plate excellent in the delayed fracture resistance of the shear surface characterized by performing said winding at 500 degreeC or more.
Record
C− (12/51) × (V− (51/14) × N) ≧ 0.013 (1)
Mn / {V− (51/14) × N} ≦ 2.0 (2)
Where C, V, N, Mn: Content of each element (% by mass) In addition to the composition of the steel material, further, in mass%, Ti: 0.005-0.600%, Nb: 0.005-0.600%, Mo: 0.005-0.600%, Ta: 0.005-0.600%, W: 0.005-0.600% 1 type or 2 types or more selected from the above, C, Mn, N, and V are replaced by the above formula (1), and the following formulas (3) to (5) according to the relationship of Ti, N, and S 9. The method according to claim 8, further comprising adjusting any one of the formula (2) in place of the formula (2) so as to satisfy the following formula (6) or (7) according to the relationship between Ti and N. A method for producing a high-strength thin steel sheet as described in 1.
Record
When Ti− (48/14) × N− (48/32) × S ≧ 0,
C− (12/51) × V− (12/48) × (Ti− (48/14) × N− (48/32) × S) − (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (3)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N ≧ 0,
C− (12/51) × V− (12/93) × Nb− (12/96) × Mo− (12/181) × Ta− (12/184) × W ≧ 0.013 (4)
When Ti− (48/14) × N− (48/32) × S <0 and Ti− (48/14) × N <0,
C− (12/51) × (V− (51/14) × (N− (14/48) × Ti)) − (12/93) × Nb− (12/96) × Mo− (12/181 ) × Ta− (12/184) × W ≧ 0.013 (5)
When Ti− (48/14) × N ≧ 0, Mn / V ≦ 2.0 (6)
When Ti− (48/14) × N <0,
Mn / (V− (51/14) × (N− (14/48) × Ti)) ≦ 2.0 (7)
Here, C, V, N, Mn, Ti, S, Nb, Mo, Ta, W: Content of each element (mass%)   The method for producing a high-strength thin steel sheet according to claim 8 or 9, further comprising B: 0.0003 to 0.0050% by mass% in addition to the composition of the steel material.   In addition to the composition of the steel material, the steel material further contains one or more selected from Cr: 0.01 to 1.0%, Ni: 0.01 to 1.0%, and Cu: 0.01 to 1.0% by mass%. The method for producing a high strength thin steel sheet according to any one of claims 8 to 10.   The method for producing a high-strength thin steel sheet according to any one of claims 8 to 11, further comprising Sb: 0.005 to 0.050% by mass% in addition to the composition of the steel material.   In addition to the composition of the steel material, the steel material further contains one or more selected from Ca: 0.0005 to 0.01% and REM: 0.0005 to 0.01% by mass%. A method for producing a high-strength thin steel sheet according to any one of 8 to 12. When performing the pickling and annealing on the hot-rolled thin steel sheet after the winding,
In the annealing, the soaking temperature is set to a temperature in the range of 750 ° C. or less, and the temperature range from 500 ° C. to the soaking temperature is heated at an average heating rate of 1 ° C./s or more, soaking at the soaking temperature. Soaking time as 300 s or less, and after soaking, cooling to 500 ° C. at a cooling rate of 1 ° C./s or more,
The method for producing a high-strength thin steel sheet according to any one of claims 8 to 13, wherein a plating treatment is performed in which the plating is immersed in a zinc plating bath having a bath temperature of 420 to 500 ° C in the course of the cooling of the annealing.   The high-strength thin steel sheet according to claim 14, wherein after the plating treatment, the heating temperature is further reheated to 460 to 600 ° C., and an alloying treatment of a plating layer that is held at the heating temperature for 1 s or more is performed. Production method.   The method for producing a high-strength thin steel sheet according to any one of claims 8 to 13, wherein after the winding, processing of a sheet thickness reduction rate: 0.1 to 3.0% is further applied.   The method for producing a high-strength thin steel sheet according to claim 14, wherein after the plating treatment, processing of a plate thickness reduction rate: 0.1 to 3.0% is applied.   The method for producing a high-strength thin steel sheet according to claim 15, wherein after the alloying treatment, processing of a sheet thickness reduction rate of 0.1 to 3.0% is further applied.

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