JP2005154906A - High tensile strength cold-rolled steel sheet having excellent strain age hardening property, and its production method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high tensile strength cold rolled steel sheet having excellent strain age hardening properties, and suitable as an automobile body, and to provide its production method. <P>SOLUTION: In the method of producing a high tensile strength cold-rolled steel sheet, a slab having a composition containing ≤0.02% Al and 0.0050 to 0.0250% N and in which N/Al is controlled to ≥0.3 is subjected to hot rolling at an FDT (finishing mill delivery temperature) of ≥800°C, and is coiled at a CT (coiling temperature) of ≤750°C. Then, after cold rolling, the steel is subjected to continuous annealing at a recrystallization temperature to 900°C, is subjected to primary cooling so as to be rapidly cooled to ≤500°C and is subjected to secondary cooling in which residence time in the temperature range of the primary cooling stopping temperature to ≥400°C is controlled to ≤300 s to form a steel sheet having a structure including an F phase with a grain size of ≤10 μm by ≥50% and N in a solid solution state by ≥0.0010%. In the method, it is also possible that the continuous annealing temperature is controlled to the two phase region of Ac<SB>1</SB>to Ac<SB>3</SB>, and cooling where the average cooling rate from 600 to 300°C is controlled to CR defined in accordance with the contents of the alloy elements or above, thus the steel is provided with a structure comprising a F phase by ≥50% and a M phase by ≥3%. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  TECHNICAL FIELD The present invention relates to a high workability high-tensile cold-rolled steel sheet that is suitable mainly for automobile bodies, and in particular, a high-tensile cold-rolled steel sheet having a tensile strength (TS) of 440 MPa or more and excellent strain aging hardening characteristics, and a method for producing the same. About. The high-tensile cold-rolled steel sheet of the present invention is used in a wide range of applications, from those used for relatively light processing such as being formed into pipes by mild bending and roll forming to those used for relatively severe drawing. It is suitable for. In addition, the steel plate in this invention shall include a steel plate and a steel strip.

  In the present invention, “excellent strain age hardening characteristics” means that after pre-deformation with a tensile strain of 5%, when subjected to aging treatment at a temperature of 170 ° C. for 20 minutes, the deformation stress increases before and after this aging treatment. The amount (denoted as BH amount; BH amount = (yield stress after aging treatment) − (predeformation stress before aging treatment)) is 80 MPa or more and before and after strain aging treatment (predeformation + the aging treatment) It means that the amount of increase in tensile strength (denoted as ΔTS; ΔTS = (tensile strength after aging treatment) − (tensile strength before pre-deformation)) is 40 MPa or more.

In connection with recent exhaust gas regulations due to global environmental problems, the reduction of vehicle weight in automobiles has become an extremely important issue. In order to reduce the weight of an automobile body, it is effective to increase the strength of a steel plate used in large quantities, that is, to apply a high-tensile steel plate to reduce the thickness of the steel plate to be used.
However, even automobile parts that use thin high-strength steel sheets must exhibit the necessary and sufficient performance according to their roles. Such performance includes, for example, static strength against bending and torsional deformation, fatigue resistance, and impact resistance. Therefore, a high-tensile steel plate applied to automobile parts needs to have excellent properties after forming.

Also, in the process of making automotive parts, press forming is performed on the steel sheet, but if the strength of the steel sheet is too high,
(I) The shape freezing property decreases,
(B) Since the ductility is reduced, problems such as cracking and necking occur during molding.
As a result, the application of high-tensile steel sheets to automobile bodies has been hindered.

  As a technique for overcoming this, for example, in the case of cold-rolled steel sheets for outer panel panels, steel sheets are known in which ultra-low carbon steel is used as a raw material, and finally the amount of C remaining in a solid solution state is controlled within an appropriate range. Yes. This kind of steel sheet is kept soft during press forming, ensuring shape freezing and ductility, and yield stress using the strain age hardening phenomenon that occurs in the paint baking process of about 170 ℃ × 20 min. It is intended to obtain a rise and secure dent resistance. In this type of steel plate, C is dissolved and soft in the steel at the time of press forming. On the other hand, after press forming, the solid solution C adheres to dislocations introduced at the time of press forming in the paint baking process, yielding. Stress increases.

However, in this type of steel sheet, the yield stress increase due to strain age hardening is kept low from the viewpoint of preventing the occurrence of the strainer strain that becomes a surface defect. For this reason, the place which actually contributes to the weight reduction of components is small.
That is, in order to reduce the weight of a part, it is necessary not only to increase the yield stress due to strain aging but also to increase the strength characteristics when the deformation progresses further. In other words, an increase in tensile strength after strain aging is desired.

On the other hand, for applications where the appearance is not a problem, the bake hardenability can be further improved by using a solid solution N to further increase the amount of bake-hardening and making the structure a composite structure of ferrite and martensite. A further improved steel sheet has been proposed.
For example, Patent Document 1 discloses a steel containing C: 0.02 to 0.15%, Mn: 0.8 to 3.5%, P: 0.02 to 0.15%, Al: 0.10% or less, and N: 0.005 to 0.025% at a temperature of 550 ° C or less. Discloses a method for producing a high-strength thin steel sheet having good ductility and spot weldability, both of which are hot-rolling wound on the steel and controlled annealing heat treatment after annealing after cold rolling. The steel sheet manufactured by the technique described in Patent Document 1 has a mixed structure composed of a low-temperature transformation product phase mainly composed of ferrite and martensite, and is excellent in ductility, and is baked by coating with positively added N. It is intended to obtain a high strength by utilizing the strain aging at the time.

However, in the technique described in Patent Document 1, mechanical properties such as an increase in yield stress YS due to strain age hardening is large but an increase in tensile strength TS is small, and an increase in yield stress YS varies greatly. Therefore, the steel sheet cannot be expected to be thin enough to contribute to the weight reduction of automobile parts currently required.
Patent Document 2 discloses a uniform bainite having a component composition including C: 0.08 to 0.20%, Mn: 1.5 to 3.5% and the balance Fe and unavoidable impurities and having a ferrite content of 5% or less. A bake-hardening high-tensile cold-rolled thin steel plate composed of bainite containing part martensite is disclosed. The cold-rolled steel sheet described in Patent Document 2 is not conventionally used by rapidly cooling the temperature range of 400 to 200 ° C. in the cooling process after continuous annealing and then gradually cooling the structure to make the structure mainly composed of bainite. It is intended to obtain a high bake hardening amount.

  However, in the steel sheet described in Patent Document 2, the yield strength is increased after baking and a high bake hardening amount that has not been obtained in the past can be obtained, but it cannot be increased by the tensile strength and applied to a strength member. In this case, improvement in fatigue resistance and impact resistance after molding cannot be expected. For this reason, the problem that it cannot apply to the use for which fatigue resistance, impact resistance, etc. are requested | required strongly remained.

Further, a steel sheet that is subjected to heat treatment after press forming to increase not only the yield stress but also the tensile strength is proposed, although it is a hot-rolled steel sheet.
For example, in Patent Document 3, steel containing C: 0.02 to 0.13%, Si: 2.0% or less, Mn: 0.6 to 2.5%, sol.Al: 0.10% or less, N: 0.0080 to 0.0250%, 1100 ° C. or more To 850 to 900 ℃, and then finish rolling at 850 to 900 ℃, then cool down to a temperature of less than 150 ℃ at a cooling rate of 15 ℃ / s and wind up, mainly ferrite and martensite A method for producing a hot-rolled steel sheet having a composite structure has been proposed. However, in the steel sheet manufactured by the technique described in Patent Document 3, although tensile strength increases with yield stress due to strain aging hardening, it is wound at an extremely low winding temperature of less than 150 ° C. There was a problem that was large. In addition, there is a large variation in the amount of increase in yield stress after press molding-paint baking treatment, and there is also a problem that the hole expansion rate (λ) is low, stretch flangeability is lowered, and press formability is insufficient.

  Moreover, as a high strength steel plate having a relatively high yield stress, there is a so-called precipitation strengthened steel in which carbonitride forming elements such as Ti, Nb, and V are added and strengthened by these fine precipitates. Regardless of hot-rolled steel sheets that have undergone a process of sufficiently retaining heat after coiling, cold-rolled steel sheets are difficult to sufficiently precipitate in a short-time continuous annealing process, and have a high yield ratio (tensile strength). It was difficult to produce a steel sheet having a yield stress ratio to YS / TS). In particular, if the C content is reduced in consideration of weldability, there is a problem that it becomes more difficult to obtain a high yield ratio because the amount of precipitates itself decreases in a region where the C content is low.

Furthermore, although the above-described conventional steel plate is excellent in strength evaluation after the baking treatment by a simple tensile test, there is a large variation in strength when plastically deformed according to actual press conditions, and reliability is high. It was not always sufficient to apply to the required parts.
JP-A-60-52528 Japanese Patent Publication No. 5-24979 Japanese Patent Publication No. 8-23048

  The present invention breaks the limitations of the prior art described above, has high moldability, or even higher impact resistance and stable quality characteristics, and has sufficient strength as an automobile part after being molded into an automobile part. To provide a high-tensile cold-rolled steel sheet excellent in strain age hardening characteristics that can sufficiently contribute to weight reduction of the obtained automobile body, and a manufacturing method that can manufacture these steel sheets industrially at low cost and without disturbing the shape. Objective. The strain age hardening characteristics in the present invention are targeted at a BH amount of 80 MPa or more and a ΔTS of 40 MPa or more under the aging conditions of holding at a temperature of 170 ° C. for 20 minutes after pre-deformation with a tensile strain of 5%.

  In order to achieve the above-mentioned problems, the present inventors manufactured steel sheets with various compositions and manufacturing conditions, and conducted many material evaluation experiments. As a result, in the field where high workability is required, N, which has not been actively used so far, is used as a strengthening element, and the large strain age hardening phenomenon expressed by the action of this strengthening element is advantageously utilized. It has been found that it is possible to easily achieve both improvement in moldability and increase in strength after molding.

  Furthermore, in order to make the best use of the strain age hardening phenomenon caused by N, the present inventors advantageously combined the strain age hardening phenomenon caused by N with the paint baking conditions of automobiles or more actively with the heat treatment conditions after molding. For that purpose, it was found that it is effective to control the microscopic structure of steel sheet and the amount of solute N within a certain range by optimizing hot rolling conditions, cold rolling and cold rolling annealing conditions. . In addition, in order to stably develop the strain age hardening phenomenon due to N, it has been found that it is important to control the Al content according to the N content particularly in terms of composition. In addition, the present inventors have made full use of N without the problem of room temperature aging degradation, which has been a problem in the past, by making the microstructure of the steel sheet the main phase of ferrite and the average grain size to be 10 μm or less. I found what I could do.

Furthermore, the present inventors can achieve a low yield ratio by making the microstructure of the steel sheet a ferrite main phase and a structure containing a martensite phase in an area ratio of 3% or more as the second phase, It has been found that the ductility and workability are improved, and the strain age hardening phenomenon expressed by N is advantageously used to increase the strength after processing and to improve the impact resistance as part characteristics.
That is, the present inventors use N as a reinforcing element, control the Al content to an appropriate range according to the N content, optimize the hot rolling conditions, cold rolling and cold rolling annealing conditions, By optimizing the visual structure and solid solution N, the formability that is much superior to conventional solid solution strengthened C-Mn steel plates and precipitation strengthened steel plates, and strain aging not found in the above conventional steel plates It has been found that a steel sheet having hardening characteristics or excellent impact resistance characteristics as part characteristics can be obtained.

  In addition, the steel sheet of the present invention is higher in strength after paint baking treatment by a simple tensile test than the conventional steel sheet, and further, there is little variation in strength when plastically deformed according to actual press conditions, and stable component strength Characteristics are obtained. For example, a portion where the strain is greatly applied and the plate thickness is reduced has a larger curing allowance than the other portions, and it is in a direction to equalize when evaluated with a load capacity of (plate thickness) x (strength), and the strength as a part is stable To do. On the other hand, the strength increase due to strain aging is small in the portion where the strain is small, but the necessary strength can be secured because the original strength is high, so the strength stability is improved in actual parts formed with various ranges of strain. Contributes favorably.

The present invention has been completed with further studies based on the above findings. That is, the gist of the high-tensile cold-rolled steel sheet according to the first invention is as follows.
(1) By mass%, Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al is 0.3 or more, and solid solution N contains 0.0010% or more, and average crystal grain size is 10 μm or less A high-tensile cold-rolled steel sheet having a tensile strength of 440 MPa or more and excellent strain age hardening characteristics, characterized by having a structure containing 50% or more of the above ferrite phase.
(2) The composition is% by mass, C: 0.15% or less, Si: 2.0% or less, Mn: 3.0% or less, P: 0.08% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050 -0.0250%, N / Al is 0.3 or more, solid solution N is contained 0.0010% or more, and the balance is composed of Fe and inevitable impurities. Tensile cold-rolled steel sheet (3) The high-tensile cold-rolled steel sheet according to (2), further comprising one group or two or more groups of the following groups a to d in mass% in addition to the above composition.

Group a: One or more of Cu, Ni, Cr, and Mo total 1.0% or less in total Group b: One or two or more of Nb, Ti, and V total 0.1% or less in total c Group: B 0.0030% or less d group: One or two of Ca and REM in total 0.0010 to 0.010%
(4) The composition contains, by mass%, C: 0.15% or less, Mn: 3.0% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050 to 0.0250%, and Mo: 0.05 to 1.0%, Cr: One or two of 0.05 to 1.0%, N / Al is 0.3 or more, N in solid solution is 0.0010% or more, and the balance consists of Fe and inevitable impurities (1), wherein the structure includes a ferrite phase having an average crystal grain size of 10 μm or less in an area ratio of 50% or more, and further includes a martensite phase in an area ratio of 3% or more. The high-tensile cold-rolled steel sheet described.
(5) The high-tensile cold-rolled steel sheet according to (4), further including one group or two or more groups among the following j groups to m groups in addition to the composition.

Group j: Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01% or two or more k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: One or more of 0.01 to 0.2% Group 1: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 or 2 types m group: 1 or 2 types of Ca and REM in total 0.0010 ~ 0.010%
(6) The high-tensile cold-rolled steel sheet according to any one of (1) to (5), wherein the high-tensile cold-rolled steel sheet has a thickness of 3.2 mm or less.
(7) A high-tensile cold-rolled steel sheet obtained by electroplating or hot-plating the high-tensile cold-rolled steel sheet according to any one of (1) to (6).

Moreover, the summary of the manufacturing method of the high-tensile cold-rolled steel sheet according to the second invention is as follows.
That is, according to the second aspect of the present invention, a steel slab having a composition containing, by mass%, Al: 0.02% or less, N: 0.0050 to 0.0250%, and N / Al of 0.3 or more is used. Heating to the above, rough rolling into a sheet bar, finishing rolling to a finish rolling exit temperature of 800 ° C. or higher, and rolling up to 750 ° C. A rolling process, a cold rolling process in which the hot-rolled sheet is pickled and cold-rolled to form a cold-rolled sheet, and the cold-rolled sheet is annealed at a predetermined temperature for a predetermined time, and then cooled to a predetermined temperature. A method for producing a high-tensile cold-rolled steel sheet having a tensile strength of 440 MPa or more and excellent strain aging hardening characteristics, characterized by sequentially performing a cold-rolled sheet annealing step of cooling at a speed.

In the second aspect of the present invention, C: 0.15% or less, Si: 2.0% or less, Mn: 3.0% or less, P: 0.08% or less, S: 0.02% or less, Al: 0.02% or less, N : 0.0050 to 0.0250% is included, and N / Al is 0.3 or more, or further a group a to d group a group: one or more of Cu, Ni, Cr, and Mo is 1.0% or less in total b Group: One or more of Nb, Ti, and V in total 0.1% or less c Group: B in 0.0030% or less d Group: One or two of Ca and REM in total 0.0010 to 0.010%
A steel slab containing one group or two or more groups selected from the above, preferably having a composition comprising the balance Fe and unavoidable impurities, is heated to a slab heating temperature of 1000 ° C. or higher, and is roughly rolled into a sheet bar. The sheet bar is subjected to finish rolling at a finish rolling exit temperature of 800 ° C. or higher, and after finish rolling, cooling is preferably started within 0.5 seconds, followed by rapid cooling at a cooling rate of 40 ° C./s or more. : Hot rolling step of 750 ° C. or lower, preferably winding temperature: 650 ° C. or lower to obtain a hot rolled sheet, and cold rolling step of pickling and cold rolling the hot rolled sheet to obtain a cold rolled sheet And the cold-rolled sheet is annealed at a recrystallization temperature of 900 ° C. or lower and a holding time of 10 to 60 s, and then cooled to a temperature range of 500 ° C. or lower at a cooling rate of 10 to 300 ° C./s. Cooling, and then the residence time in the temperature range of 400 ° C or higher below the primary cooling stop temperature It is preferable to sequentially perform a cold-rolled sheet annealing step in which secondary cooling is performed at a rate of 300 s or less. In the second aspect of the present invention, the elongation ratio: 1.0 is further added following the cold-rolled plate annealing step. It is preferable to apply -15% temper rolling or leveler processing. In the second aspect of the present invention, it is preferable to join adjacent sheet bars between the rough rolling and the finish rolling. In the second aspect of the present invention, the rough rolling and the finish rolling are performed. It is preferable to use either or both of a sheet bar edge heater for heating the width end portion of the seat bar and a sheet bar heater for heating the length end portion of the seat bar.

Further, in the second aspect of the present invention, by mass%, C: 0.15% or less, Mn: 3.0% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050 to 0.0250%, and Mo: 0.05 to 1.0%, Cr: 1 or 2 of 0.05 to 1.0%, and N / Al is 0.3 or more, or further j group to m group j group: Si: 0.05 to 1.5 %, P: 0.03 to 0.15%, B: 0.0003 to 0.01%, 1 type or 2 types or more k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: 0.01 to 0.2%, 1 type or 2 Species or more Group I: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 or 2 types m Group: 1 or 2 types of Ca and REM in total 0.0010 to 0.010%
A steel slab containing one group or two or more groups selected from the above, preferably having a composition comprising the balance Fe and unavoidable impurities, is heated to a slab heating temperature of 1000 ° C. or higher, and is roughly rolled into a sheet bar. The sheet bar is subjected to finish rolling at a finish rolling exit temperature of 800 ° C. or higher, and after finish rolling, cooling is preferably started within 0.5 seconds, and the cooling rate is rapidly cooled at 40 ° C./s or more, and the coiling temperature : A hot rolling process for winding a hot rolled sheet at 750 ° C. or lower, a cold rolling process for pickling and cold rolling the hot rolled sheet to form a cold rolled sheet, and (Ac Annealing is performed at a temperature of ₁ transformation point) to (Ac ₃ transformation point) and a holding time of 10 to 120 s, and then the average cooling rate between 600 to 300 ° C. is set as follows (1) or (2) Formula B <0.0003% in the case of
log CR = -1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.95 (1)
When B ≧ 0.0003%
log CR = −1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.40 …… (2)
(Here, CR: cooling rate (° C./s), Mn, Mo, Cr, Si, P, Cu, Ni: content of each element (mass%))
It is preferable to perform a cold-rolled sheet annealing step for cooling at a critical cooling rate CR defined by (1) or more in order to obtain a cold-rolled steel sheet having excellent workability and impact resistance. In the second aspect of the present invention, after the cold-rolled sheet annealing step, temper rolling or leveler processing with an elongation of 1.0 to 15% is preferably performed. In the second aspect of the present invention, it is preferable to join adjacent sheet bars between the rough rolling and the finish rolling. In the second aspect of the present invention, the rough rolling and the finish rolling are performed. It is preferable to use either or both of a sheet bar edge heater for heating the width end portion of the seat bar and a sheet bar heater for heating the length end portion of the seat bar.

  According to the present invention, the yield stress is 80 MPa or more and the tensile strength is 40 MPa or more due to the pre-deformation-paint baking process, and both high strain age hardening properties and high formability, or even higher impact resistance properties are increased. A high-tensile cold-rolled steel sheet can be manufactured at low cost without disturbing the shape, and has a remarkable industrial effect. Furthermore, when the high-tensile cold-rolled steel sheet of the present invention is applied to automobile parts, it is possible to obtain stable and high part characteristics by increasing the tensile strength as well as the yield stress by paint baking treatment, etc. For example, from 2.0 mm to 1.6 mm thick, it is possible to reduce the grade by 1 grade from the conventional level, and there is also an effect that it can sufficiently contribute to weight reduction of the car body.

The steel sheet of the present invention is a high-tensile cold-rolled steel sheet having a tensile strength of 440 MPa or more and excellent strain age hardening characteristics. First, the reasons for limiting the composition of the steel sheet of the present invention will be described. Hereinafter, the mass% is simply referred to as%.
The steel sheet of the present invention has a composition containing Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al being 0.3 or more, and solid solution N being 0.0010% or more.

Al: 0.02% or less
Al is an element that acts as a deoxidizer and is effective in improving the cleanliness of the steel, and is also an element that refines the structure of the steel sheet. In the present invention, Al is preferably contained in an amount of 0.001% or more. On the other hand, excessive Al content deteriorates the surface properties of the steel sheet, further reduces the solid solution N, which is an important constituent element of the present invention, and causes a shortage of solid solution N that contributes to the strain age hardening phenomenon. When the conditions vary, the strain age hardening characteristic, which is a feature of the present invention, tends to vary. For this reason, in the present invention, the Al content is limited to as low as 0.02% or less. From the viewpoint of material stability, Al is preferably 0.015% or less.

N: 0.0050-0.0250%
N is an element that increases the strength of the steel sheet by solid solution strengthening and strain age hardening, and is the most important element in the present invention. N also has a function of lowering the transformation point of steel, and the content of N is also useful for stabilizing the operation in a situation where rolling with a thin material that greatly interrupts the transformation point is avoided. In the present invention, an appropriate amount of N is contained and the production conditions are controlled, thereby securing a necessary and sufficient amount of N in a solid solution state in a cold-rolled product or a plated product, thereby strengthening solid solution and strain aging. Strength of steel (YS, TS) increase effect in hardening, TS440MPa or more, bake hardening amount (BH amount) 80MPa or more, increase in tensile strength before and after strain aging treatment ΔTS40MPa or more It is possible to stably satisfy the physical property requirements.

  When N is less than 0.0050%, the above-described strength increasing effect is not likely to appear stably. On the other hand, if N exceeds 0.0250%, the rate of occurrence of internal defects in the steel sheet increases, and slab cracking during continuous casting occurs frequently. For this reason, N was made into the range of 0.0050-0.0250%. Note that N is more preferably in the range of 0.0070 to 0.0170% from the viewpoint of improving the stability and yield of the material considering the entire manufacturing process. If the N amount is within the range of the present invention, there is no adverse effect on weldability such as spot welding and arc welding.

Solid solution state N: 0.0010% or more In order to ensure sufficient strength in cold-rolled products and to fully exhibit strain age hardening due to N, solid solution state N in steel (also called solid solution N) Must be present in an amount (concentration) of at least 0.0010%.
Here, the solute N amount is obtained by subtracting the precipitated N amount from the total N amount in the steel. As an analysis method for the amount of precipitated N, it is effective to obtain by an electrolytic extraction analysis method using a constant potential electrolysis method according to the results of a comparative study of various analysis methods by the present inventors. In addition, there are an acid decomposition method, a halogen method, and an electrolytic method as a method for dissolving the base iron used for the extraction analysis. Among these, the electrolytic method can dissolve only the iron core stably without decomposing extremely unstable precipitates such as carbides and nitrides. Electrolysis is performed at a constant potential using an acetylacetone system as the electrolytic solution. In the present invention, the result of measuring the amount of precipitated N using a constant potential electrolysis method showed the best correspondence with the actual component strength.

For this reason, in the present invention, the residue extracted by the constant potential electrolysis method is chemically analyzed to determine the amount of N in the residue, which is used as the amount of precipitated N.
In order to obtain a higher BH amount and ΔTS, the solid solution N amount is 0.0020% or more. To obtain a higher value, 0.0030% or more. To obtain a higher value, 0.0050% or more. It is preferable to do this.

N / Al (ratio of N content to Al content): 0.3 or more In the product state, in order to make solid solution N stable and remain at 0.0010% or more, the amount of Al, which is an element that strongly fixes N, is changed. Need to be restricted. As a result of examining a steel sheet in which the combination of N content and Al content within the composition range of the present invention is changed over a wide range, in order to make the solid solution N in cold-rolled products and plated products 0.0010% or more, the Al content is It was found that N / Al needs to be 0.3 or more when limited to 0.02% or less. That is, the Al content is limited to (N content) /0.3 or less.

Furthermore, the steel sheet of the present invention has a structure containing a ferrite phase having an average crystal grain size of 10 μm or less in an area ratio of 50% or more. The structure of the steel sheet of the present invention will be described.
Area ratio of ferrite phase: 50% or more The cold-rolled steel sheet of the present invention is intended for use in automobile steel sheets and the like that require high workability, and in order to ensure ductility, the ferrite phase is 50% in area ratio. % Or more of the ferrite phase as the main phase. If the area ratio of the ferrite phase is less than 50%, it becomes difficult to ensure the ductility necessary for a steel sheet for automobiles that requires high workability. In addition, when better ductility is required, the area ratio of the ferrite phase is preferably 75% or more. In addition, the ferrite referred to in the present invention includes not only the ordinary meaning of ferrite (polygonal ferrite) but also bainitic ferrite and acicular ferrite not containing carbide.

The second phase other than the ferrite phase is not particularly limited, but is preferably a single phase or a mixed phase of bainite or martensite from the viewpoint of increasing the strength. Moreover, it is good also as a structure | tissue which mainly has pearlite as a 2nd phase. In addition, the retained austenite may appear in an area ratio of less than about 3% within the composition range of the steel sheet of the present invention.
Average crystal grain size of ferrite phase: 10 μm or less In the present invention, as the crystal grain size, the value calculated by the quadrature method prescribed in ASTM from the sectional structure photograph and the nominal grain obtained by the cutting method prescribed by ASTM from the sectional structure photograph The larger one of the diameters (for example, see Umemoto et al .: Heat treatment, 24 (1984), 334) is used.

  The cold-rolled steel sheet of the present invention secures a predetermined amount of solute N as a product, but according to the results of experiments and examinations by the present inventors, the average of the ferrite phase is maintained even if the amount of solute N is kept constant. It has been found that when the crystal grain size exceeds 10 μm, the strain age hardening characteristics vary greatly. In addition, the deterioration of mechanical properties when stored at room temperature becomes significant. The details of this mechanism are currently unknown, but one of the causes of the variation in strain age hardening characteristics is the crystal grain size. Segregation and precipitation of alloy elements at the grain boundaries, as well as the processing and heat treatments affecting them. It is presumed to be related to the effects of Therefore, in order to stabilize the strain age hardening characteristics, the average crystal grain size of the ferrite phase needs to be 10 μm or less. In order to stably obtain a further increase in the BH amount and ΔTS amount, the average crystal grain size of ferrite is preferably 8 μm or less.

The cold-rolled steel sheet of the present invention having the composition and structure described above is a cold-rolled steel sheet having a tensile strength TS of 440 MPa or more and excellent strain age hardening characteristics.
A steel sheet having a TS of less than 440 MPa cannot be widely applied to members having structural members. In order to further expand the applicable range, it is desirable that TS is 500 MPa or more.
In the present invention, “excellent in strain age hardening characteristics” means that, as described above, after pre-deformation with a tensile strain of 5%, when subjected to aging treatment at a temperature of 170 ° C. for 20 minutes, before and after this aging treatment. Tensile stress before and after strain aging treatment (pre-deformation + aging treatment) when deformation stress increase (BH amount; BH amount = yield stress after aging treatment-pre-deformation stress before aging treatment) is 80 MPa or more It means that the amount of increase in strength (denoted as ΔTS; ΔTS = tensile strength after aging treatment−tensile strength before pre-deformation) is 40 MPa or more.

  When the strain age hardening characteristic is specified, the amount of pre-strain (pre-deformation) is an important factor. Assuming the deformation mode applied to the steel sheet for automobiles, the present inventors investigated the influence of the pre-strain amount on the strain age hardening characteristics. As a result, (i) the deformation stress in the deformation mode is extremely high. Except in the case of deep drawing, the strain can be roughly organized by the amount equivalent to uniaxial strain (tensile strain). (Ii) In actual parts, the amount of strain equivalent to uniaxially exceeds 5%. It was found that it corresponds well with the strength (YS and TS) obtained after the strain aging treatment with 5% pre-strain. Based on this knowledge, in the present invention, the pre-deformation of the strain aging treatment is set to 5% tensile strain.

  Conventional coating baking conditions of 170 ° C x 20 min have been adopted as standard. In general, in order to earn a hardened amount, it is advantageous to hold at a higher temperature for a longer time unless softened by excessive aging. When strain of 5% or more is applied, curing can be achieved even by a more gradual (low temperature side) treatment, in other words, a wider range of aging conditions can be taken.

  Specifically, in the steel sheet of the present invention, the lower limit of the heating temperature at which hardening becomes significant after pre-deformation is approximately 100 ° C. On the other hand, when the heating temperature exceeds 300 ° C., the curing reaches its peak, and on the contrary, there is a tendency to slightly soften, and the occurrence of thermal distortion and temper color becomes conspicuous. As for the holding time, when the heating temperature is about 200 ° C., if it is about 30 seconds or longer, substantially sufficient curing can be achieved. In order to obtain larger and more stable curing, it is preferable that the holding time is 60 seconds or longer. However, if the holding time exceeds 20 minutes, further curing cannot be expected, and the production efficiency is significantly reduced, which is disadvantageous in practical use.

  From the above, in the present invention, it was determined that the aging treatment conditions were evaluated at 170 ° C. which is the heating temperature of the conventional paint baking treatment conditions and the holding time was 20 minutes. Even with the low temperature heating and short-time aging treatment conditions in which sufficient hardening cannot be achieved with conventional paint-baked steel sheets, large hardening is stably achieved with the steel sheets of the present invention. The heating method is not particularly limited, and any of induction heating, heating with a non-oxidizing flame, laser, plasma, etc., for example, can be preferably used in addition to atmospheric heating with a furnace employed for ordinary paint baking.

  The strength of parts for automobiles must be able to withstand complex stress loads from the outside. Therefore, in a steel plate, not only strength characteristics in a small strain range but also strength characteristics in a large strain range are important. In view of this point, the present inventors set the BH amount of the steel sheet of the present invention to be a material for automobile parts to 80 MPa or more and the ΔTS amount to 40 MPa or more. More preferably, the BH amount is 100 MPa or more and ΔTS50 MPa or more. In order to increase the BH amount and the ΔTS amount, the heating temperature during the aging treatment may be set to a higher temperature side and / or the holding time may be set to a longer time side.

In addition, the steel sheet of the present invention has a conventional structure in which it is extremely difficult to cause aging deterioration (a phenomenon in which YS increases and El (elongation) decreases) even if it is left for about a year at room temperature when not formed. Has no advantages.
By the way, the effect of the present invention can be exhibited even when the product plate thickness is relatively thick. However, when the product plate thickness exceeds 3.2 mm, a necessary and sufficient cooling rate can be secured in the cold-rolled plate annealing process. However, strain aging occurs during continuous annealing, making it difficult to obtain the target strain age hardening characteristics as a product. Therefore, the thickness of the steel sheet of the present invention is preferably 3.2 mm or less.

  In the present invention, there is no problem even if the surface of the cold-rolled steel sheet of the present invention is electroplated or hot-plated. These plated steel sheets also exhibit the same amount of TS, BH, and ΔTS as before plating. As the kind of plating, any of electrogalvanizing, hot dip galvanizing, alloying hot dip galvanizing, electrotin plating, electrochromic plating, electronickel plating, etc. can be preferably applied.

In the steel sheet of the present invention, chemical components other than the above-described Al and N can be appropriately selected according to required characteristics.
The steel sheet according to the first preferred embodiment includes the above-described composition containing Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al is 0.3 or more, and N in a solid solution state is 0.0010% or more. In addition, C: 0.15% or less, Si: 2.0% or less, Mn: 3.0% or less, P: 0.08% or less, S: 0.02% or less, or further, the following a group to d group a group: Cu, 1 type or 2 types or more of Ni, Cr and Mo in total 1.0% or less b group: 1 type or 2 types or more of Nb, Ti and V in total 0.1% or less c group: B is 0.0030% or less d group: One or two of Ca and REM in total 0.0010 to 0.010%
1 or 2 or more, and the balance is composed of Fe and inevitable impurities, and a structure including a ferrite phase having an average crystal grain size of 10 μm or less in an area ratio of 50% or more.

The reasons for limiting the composition other than Al and N of the steel sheet as the first preferred embodiment will be described. The reason for limiting the organization is the same as described above.
C: 0.15% or less C is an element that increases the strength of the steel sheet, and in order to achieve an average grain size of ferrite of 10 μm or less, which is an important constituent element of the present invention, from the viewpoint of further securing a desired strength. , 0.005% or more is preferable. More preferably, it is 0.03% or more. On the other hand, if it exceeds 0.15%, the carbide fraction in the steel sheet becomes excessive, the ductility is remarkably reduced and formability is deteriorated, and further, spot weldability, arc weldability, etc. are remarkably lowered. From the viewpoints of formability and weldability, C is limited to 0.15% or less. In addition, it is preferably 0.10% or less, and 0.08% or less for applications that require better ductility.

Si: 2.0% or less
Si is a useful element that can increase the strength of the steel sheet without significantly reducing the ductility of the steel, and is preferably contained in an amount of 0.1% or more. On the other hand, Si is an element that greatly increases the transformation point during hot rolling and makes it difficult to ensure quality and shape, or adversely affects the beauty of the steel sheet surface such as surface properties and chemical conversion properties. Then, it limited to 2.0% or less. If Si is 2.0% or less, a remarkable increase in transformation point can be suppressed by adjusting the amount of Mn added in combination, and good surface properties can be secured. In addition, when it is desired to ensure high ductility in a tensile strength TS500 MPa super high strength steel sheet, it is more preferable to contain 0.3% or more of Si from the viewpoint of balance between strength and ductility.

Mn: 3.0% or less
Mn is an effective element for preventing hot cracking due to S, and is preferably added according to the amount of S contained, and Mn is an important component of the present invention for the refinement of crystal grains. It has a great effect, and it is preferable to add it positively and use it for improving the material. From the viewpoint of stably fixing S, Mn is preferably contained in an amount of 0.2% or more.

  Mn is an element that increases the strength of the steel sheet, and is preferably contained in an amount of 1.2% or more for the strength requirement exceeding TS500 MPa. From the viewpoint of stably securing the strength, it is more preferably 1.5% or more. When the Mn content is increased to this level, variations in the mechanical properties and strain age hardening characteristics of the steel sheet with respect to fluctuations in production conditions including hot rolling conditions are reduced, which is effective in stabilizing the quality.

Further, Mn has a function of lowering the transformation point during hot rolling, and inclusion with Si can offset an increase in transformation point due to the inclusion of Si. Especially for products with thin plate thickness, quality and shape change sensitively due to changes in transformation point, so it is important to balance Mn and Si contents precisely. For these reasons, Mn / Si is more preferably 3.0 or more.
On the other hand, if Mn is contained in a large amount exceeding 3.0%, the hot deformation resistance of the steel sheet tends to increase, the spot weldability and the formability of the welded portion tend to deteriorate, and further, the formation of ferrite occurs. Since it is suppressed, the ductility tends to decrease remarkably. For this reason, Mn was limited to 3.0% or less. In applications where better corrosion resistance and formability are required, Mn is desirably 2.5% or less.

P: 0.08% or less P is an element useful as a solid solution strengthening element of steel. From this point of view, it is desirable to contain 0.001% or more. Reduce sex. Moreover, since P has a strong tendency to segregate in steel, it causes embrittlement of the weld due to it. For this reason, P was limited to 0.08% or less. In addition, when the stretch flange workability and weld toughness are particularly important, 0.04% or less is preferable. Further, it is more preferably 0.02% or less from the viewpoint of weld zone toughness.

S: 0.02% or less S is an element which exists as an inclusion in the steel sheet and causes deterioration of the ductility and further corrosion resistance of the steel sheet. In the present invention, S is limited to 0.02% or less. In applications that require particularly good workability (especially stretch flangeability and hole expansibility), the content is preferably 0.015% or less. Further, when the required level of stretch flangeability is high, S is preferably 0.008% or less. Further, in order to stably maintain the strain age hardening characteristics at a high level, the detailed mechanism is unknown, but it is preferable to reduce S to 0.008% or less.

In the first preferred embodiment of the steel sheet of the present invention, in addition to the above-described composition, the following a group to d group: a group: one or more of Cu, Ni, Cr, and Mo in total 1.0% or less b group : One or more of Nb, Ti, and V in total 0.1% or less c group: B in 0.0030% or less d group: One or two of Ca and REM in total 0.0010 to 0.010%
It is preferable to contain 1 group or 2 groups or more.

  Group a elements: Cu, Ni, Cr, and Mo are all elements that contribute to increasing the strength of the steel sheet. To obtain this effect, Cu: 0.01% or more, Ni: 0.01% or more, Cr: 0.01% As mentioned above, it is preferable to set it as Mo: 0.01% or more, It can select as needed and can contain individually or in combination. However, when the content is too large, the hot deformation resistance increases, or the chemical conversion property and the surface treatment characteristics in a broad sense are deteriorated, and the welded portion is hardened to deteriorate the weldability. For this reason, it is preferable that the elements in group a be 1.0% or less in total.

  Group b elements: Nb, Ti, and V are all elements that contribute to refinement and uniformity of crystal grains. To obtain this effect, Nb is 0.002% or more, Ti is 0.002% or more, and V: The content is preferably 0.002% or more, and can be selected alone or in combination, depending on necessity. However, when there is too much content, hot deformation resistance will increase and chemical conversion property and the surface treatment characteristic in a broad sense will deteriorate. For this reason, it is preferable that the total amount of elements in group b is 0.1% or less.

  Group c element: B is an element having an effect of improving the hardenability of the steel, and is contained as necessary for the purpose of increasing the strength of the steel by increasing the fraction of the low-temperature transformation phase other than the ferrite phase. be able to. In order to obtain this effect, B is preferably contained in an amount of 0.0002% or more. However, when the amount is too large, the hot deformability is lowered, and the solute N is reduced by generating BN. For this reason, B is preferably 0.0030% or less.

  The elements of group d: Ca and REM are all elements that are useful for controlling the form of inclusions, and are particularly preferably contained alone or in combination when there is a demand for stretch flange formability. In that case, if the total amount of elements in the d group is less than 0.0010%, the effect of controlling the shape of inclusions is insufficient. On the other hand, if it exceeds 0.010%, surface defects are conspicuous. For this reason, it is preferable to limit the elements of the d group to a total range of 0.0010 to 0.010%.

The balance other than the above components is Fe and inevitable impurities. As an inevitable impurity, O: 0.0050% or less is allowed.
The cold-rolled steel sheet of the present invention having the composition and structure described above is a cold-rolled steel sheet having a tensile strength TS of 440 MPa or more and excellent strain age hardening characteristics.
Below, the 2nd suitable aspect of this invention steel plate is demonstrated.

The steel sheet according to the second preferred embodiment contains the above-mentioned composition: Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al is 0.3 or more, and N in a solid solution state is 0.0010% or more. In addition, C: 0.15% or less, Mn: 3.0% or less, S: 0.02% or less, and Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0% Or further j group to m group j group: Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01% or two or more k group: Nb: 0.01 to 0.1% , Ti: 0.01 to 0.2%, V: 0.01 to 0.2%, 1 type or 2 types Group 1: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 type or 2 types m group: Ca, REM 1 type or 2 types in total 0.0010-0.010%
Including one or more of them, the balance consisting of Fe and inevitable impurities, and the structure contains 50% or more ferrite phase with an average crystal grain size of 10 μm or less in area ratio, and further martensite phase The steel sheet is characterized in that it has a structure containing 3% or more in terms of area ratio and is excellent in strain age-hardening properties, as well as workability and impact resistance properties.

The reasons for limiting the composition other than Al and N of the steel sheet of the second preferred embodiment will be described.
C: 0.15% or less C is an element that increases the strength of the steel sheet. Further, in order to achieve an average grain size of ferrite of 10 μm or less, which is an important component of the present invention, a desired strength is further ensured. From the viewpoint of forming a martensite phase as two phases, the content is preferably 0.005% or more. In addition, More preferably, it is 0.03% or more. On the other hand, if it exceeds 0.15%, the carbide fraction in the steel sheet becomes excessive, the ductility is remarkably lowered, the formability is deteriorated, the martensitic transformation temperature is lowered, and spot weldability and arc weldability are further remarkable. To drop. From the viewpoint of such formability, weldability, and microstructural adjustment, C is limited to 0.15% or less. In addition, it is preferably 0.10% or less, and 0.08% or less for applications that require better ductility.

Mn: 3.0% or less
Mn is an effective element for preventing hot cracking due to S, and is preferably added according to the amount of S contained, and Mn is great for crystal grain refinement, which is an important constituent of the present invention. It is effective, and it is preferable to add it positively and use it for improving the material. Further, Mn is an element that improves hardenability, and it is preferably added positively from the viewpoint of stably forming a martensite phase as the second phase. From the viewpoint of fixing S and forming a martensite phase, Mn is preferably contained in an amount of 0.2% or more.

  Mn is an element that increases the strength of the steel sheet, and is preferably contained in an amount of 1.2% or more for the strength requirement exceeding TS500 MPa. From the viewpoint of stably obtaining the strength, it is more preferably 1.5% or more. Further, when the Mn content is increased to this level, the variation in the mechanical properties and strain age hardening characteristics of the steel sheet with respect to fluctuations in manufacturing conditions including hot rolling conditions is reduced, which is effective in stabilizing the quality.

  On the other hand, if Mn is contained in a large amount exceeding 3.0%, the hot deformation resistance of the steel sheet tends to increase, the spot weldability and the formability of the welded portion tend to deteriorate, and further, the formation of ferrite occurs. Since it is suppressed, the ductility tends to decrease remarkably. For this reason, Mn was limited to 3.0% or less. In applications where better corrosion resistance and formability are required, Mn is desirably 2.5% or less.

S: 0.02% or less S is an element which exists as an inclusion in the steel sheet and causes deterioration of the ductility and further corrosion resistance of the steel sheet. In the present invention, S is limited to 0.02% or less. In applications that require particularly good workability (especially stretch flangeability and hole expansibility), the content is preferably 0.015% or less. Further, when the required level of stretch flangeability is high, S is preferably 0.008% or less. Further, in order to stably maintain the strain age hardening characteristics at a high level, the detailed mechanism is unknown, but it is preferable to reduce S to 0.008% or less.

One or two of Mo: 0.05-1.0% and Cr: 0.05-1.0%
Both Mo and Cr are elements that contribute to increasing the strength of the steel sheet, further improve the hardenability of the steel, and easily form a martensite phase as the second phase, and are contained alone or in combination. In particular, Mo and Cr have the effect of finely dispersing the martensite phase, and have the effect of easily reducing the yield strength and realizing a low yield ratio. Such an effect is recognized when the content of Mo and Cr is 0.05% or more. On the other hand, if the Mo content exceeds 1.0%, the workability and the surface treatment property are lowered, and the production cost is increased, which is economically disadvantageous. On the other hand, if the Cr content exceeds 1.0%, the plating wettability decreases. For this reason, it is preferable to limit Mo to 0.05 to 1.0% and Cr to 0.05 to 1.0%.

In the steel sheet of the second preferred embodiment of the present invention, in addition to the above composition, the following j group to m group j group: Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01% 1 type or 2 types or more of k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: 0.01 to 0.2%, 1 type or 2 types or more l group: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5% of 1 or 2 types m group: 1 or 2 types of Ca and REM in total 0.0010 to 0.010%
It is preferable to contain 1 group or 2 groups or more.

  The elements of group j: Si, P, and B are all effective elements for improving the hardenability and forming the martensite phase, and for improving the ductility and the r value to improve the workability. It is an effective element and can be contained alone or in combination as required. Such an effect is recognized in the second preferred embodiment when each Si content is 0.05% or more, P is 0.03% or more, and B is 0.0003% or more. On the other hand, if each Si content exceeds 1.5%, P content exceeds 0.15% and B content exceeds 0.01%, toughness, weldability, workability and the like deteriorate. For this reason, it is preferable to limit Si to 0.05 to 1.5%, P to 0.03 to 0.15%, and B to 0.0003 to 0.01%.

  The elements of group k: Nb, Ti, and V are all elements that contribute to the refinement and homogenization of crystal grains, and can be selected alone or in combination as required. In the second preferred embodiment, such an effect is recognized in each case of Nb: 0.01% or more, Ti: 0.01% or more, V: 0.01% or more, but each Nb: 0.1%, Ti: 0.2%, If the content exceeds V: 0.2%, the hot deformation resistance is increased, and the chemical conversion property and the surface treatment characteristics in a broad sense are deteriorated. For this reason, it is preferable to limit to Nb: 0.01-0.1%, Ti: 0.01-0.2%, V: 0.01-0.2%, respectively.

  Group l elements: Cu and Ni are elements that improve the hardenability and contribute to an increase in strength of the steel sheet, and can be selected alone or in combination as required. In the second preferred embodiment, such an effect is recognized when Cu: 0.05% or more and Ni: 0.05% or more. However, if each Cu content exceeds 1.5% and Ni exceeds 1.5%, the hot deformation resistance increases, the chemical conversion property and the surface treatment characteristics in a broad sense deteriorate, and the welded part hardens to form the welded part. Deteriorates. For this reason, it is preferable to limit to Cu: 0.05-1.5% and Ni: 0.05-1.5%, respectively.

  The elements of group m: Ca and REM are all elements that are useful for controlling the form of inclusions, and are particularly preferably contained alone or in combination when there is a demand for stretch flange formability. In that case, if the total amount of elements in the d group is less than 0.0010%, the effect of controlling the shape of inclusions is insufficient. On the other hand, if it exceeds 0.010%, surface defects are conspicuous. For this reason, it is preferable to limit the elements of the m group to a total range of 0.0010 to 0.010%.

The balance other than the above components is Fe and inevitable impurities. As an inevitable impurity, O: 0.0050% or less is acceptable.
Next, the structure of the steel sheet according to the second preferred embodiment of the present invention will be described.
The structure of the steel sheet of the second preferred embodiment of the present invention includes a ferrite phase having an average crystal grain size of 10 μm or less as a main phase in an area ratio of 50% or more, and a martensite phase as a second phase in an area ratio of 3% or more. It is a microscopic tissue.

  By including the ferrite phase in an area ratio of 50% or more, as in the first preferred embodiment, it is possible to ensure the ductility necessary for a steel sheet for automobiles that requires high workability. In addition, when better ductility is required, the area ratio of the ferrite phase is preferably 75% or more. When the area ratio of the ferrite phase exceeds 97%, the effect as a composite structure cannot be expected. The martensite phase as the second phase is dispersed mainly at the grain boundaries of the ferrite phase that is the main phase. Martensite is a hard phase and has the effect of increasing the strength of the steel sheet by strengthening the structure. Furthermore, since the generation of movable dislocations is accompanied during transformation, it has the effect of improving ductility and reducing the yield ratio of the steel sheet. These effects become significant when the martensite phase is present in an area ratio of 3% or more. In addition, when it exists exceeding 30%, there exists a problem of a ductility fall. For this reason, the martensite phase as the second phase is 3% or more, preferably 30% or less, more preferably 20% or less in terms of area ratio. As the second phase, in addition to such an amount of martensite, 10% or less of bainite may be contained without any problem.

In addition, the average crystal grain size of the ferrite phase is set to 10 μm or less in order to stabilize the strain age hardening characteristics as in the first preferred embodiment.
The steel sheet of the second preferred embodiment of the present invention having the composition and structure described above has a tensile strength TS of 440 MPa or more, excellent strain age hardening characteristics, and high tensile cold rolling excellent in workability and impact resistance characteristics. It becomes a steel plate.

Next, the manufacturing method of this invention steel plate is demonstrated.
The steel sheet of the present invention basically has a composition within the above-described range, that is, a steel slab having a composition containing Al: 0.02% or less, N: 0.0050-0.0250%, and N / Al being 0.3 or more. Slab heating temperature: heated to 1000 ° C or higher, roughly rolled into a sheet bar, finished rolling to a finish rolling exit temperature of 800 ° C or higher, and wound at a winding temperature of 750 ° C or lower A hot-rolling step for forming a hot-rolled sheet, a cold-rolling step for pickling and cold-rolling the hot-rolled plate to form a cold-rolled plate, and annealing for holding the cold-rolled plate at a predetermined temperature for a predetermined time Followed by a cold-rolled sheet annealing step for cooling at a predetermined cooling rate.

  The slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macro segregation of components, but may be produced by an ingot forming method or a thin slab continuous casting method. In addition to the normal process of once cooling the slab to room temperature and heating it again, it is rolled directly after rolling in a furnace after inserting it into a heating furnace without cooling or rolling immediately after performing a slight heat retention. Energy saving processes such as direct rolling can also be applied without problems. In particular, in order to effectively secure N in a solid solution state, direct feed rolling in which precipitation of N is delayed is one of useful techniques.

First, the reasons for limiting the hot rolling process conditions will be described.
Slab heating temperature: 1000 ° C or more The slab heating temperature is 1000 ° C in order to ensure the necessary and sufficient amount of solute N in the initial state and to meet the target value (0.0010% or more) of solute N in the product. The above is preferable. From the viewpoint of avoiding an increase in loss accompanying an increase in oxidized weight, the slab heating temperature is preferably 1280 ° C. or lower.

The slab heated under the above conditions is made into a sheet bar by rough rolling. The conditions for rough rolling need not be specified, and may be performed according to a conventional method. However, from the viewpoint of securing the solid solution N amount, it is desirable to perform the treatment in as short a time as possible.
Next, the sheet bar is finish-rolled to obtain a hot-rolled sheet.
In the present invention, it is preferable to join successive sheet bars between rough rolling and finish rolling and continuously finish rolling. As the joining means, it is preferable to use a pressure welding method, a laser welding method, an electron beam welding method or the like.

As a result, the proportion of unsteady parts (the front end and the rear end of the material to be processed) that are likely to be disturbed in finish rolling and subsequent cooling is reduced, and the stable rolling length (continuous length that can be rolled under the same conditions) is reduced. And the stable cooling length (continuous length that can be cooled with tension applied) are extended, and the shape / dimensional accuracy and yield of the product are improved.
In addition, it has become possible to easily carry out lubrication rolling for thin objects and wide widths, which was difficult to perform due to problems such as sheeting and biting by conventional single rolling for each sheet bar, reducing rolling load and roll surface pressure. This extends the life of the roll.

  In the present invention, either one or both of a sheet bar edge heater that heats the width end portion of the sheet bar and a sheet bar heater that heats the length end portion of the sheet bar is used between rough rolling and finish rolling. Thus, it is preferable to make the temperature distribution in the width direction and the longitudinal direction of the sheet bar uniform. Thereby, the material dispersion | variation in a steel plate can be made still smaller. The sheet bar edge heater and the sheet bar heater are preferably of the induction heating type.

  As for the use procedure, it is desirable to first compensate for the temperature difference in the width direction by the sheet bar edge heater. The amount of heating at this time is preferably set so that the temperature distribution range in the width direction on the finish rolling exit side is approximately 20 ° C. or less, although it depends on the steel composition and the like. Next, the temperature difference in the longitudinal direction is compensated by the sheet bar heater. The heating amount at this time is preferably set so that the length end temperature is higher by about 20 to 40 ° C. than the center temperature.

Finishing rolling exit temperature: 800 ° C. or more The finishing rolling exit temperature FDT is set to 800 ° C. or more in order to make the structure of the steel sheet uniform and fine. When the FDT is below 800 ° C., the structure becomes non-uniform, and the processed structure remains in part. Such remaining of the processed structure can be avoided by increasing the coiling temperature. However, when the coiling temperature is increased, coarse crystal grains are generated and the amount of solute N is greatly reduced, making it difficult to obtain a target tensile strength of TS440 MPa or more. In order to further improve the mechanical properties, the FDT is desirably 820 ° C. or higher. Further, from the viewpoint of preventing the occurrence of scale flaws and the like due to excessive increase in the rolling temperature, the FDT is preferably set to 1000 ° C. or less. After finish rolling, it is desirable to cool the steel sheet at an early stage in order to refine crystal grains and secure a solid solution N amount.

Cooling after finish rolling: Start cooling within 0.5 seconds after finishing rolling, rapid cooling at a cooling rate of 40 ° C / s or more In the present invention, cooling is started immediately after finishing rolling (within 0.5 seconds) The average cooling rate is desirably 40 ° C./s or more. By satisfying this condition, the high temperature region where AlN precipitates can be rapidly cooled, and N in a solid solution state can be effectively secured. If this cooling start time or cooling rate does not satisfy the above conditions, grain growth proceeds too much, making it difficult to reduce the crystal grain size, and precipitation of AlN due to strain energy introduced by rolling proceeds too much. This increases the risk that the amount of dissolved N will be deficient. From the viewpoint of ensuring the uniformity of the material and shape, the cooling rate is preferably suppressed to 300 ° C./s or less.

Winding temperature: 750 ° C. or less The steel sheet strength tends to increase as the winding temperature CT decreases. In order to ensure the target tensile strength of TS440 MPa or higher, the CT is preferably 750 ° C. or lower. If the CT is less than 200 ° C., the shape of the steel sheet tends to be disturbed, and there is a high risk of causing problems in actual operation, and the uniformity of the material tends to decrease. For this reason, it is desirable that CT be 200 ° C. or higher. When more uniform material is required, CT is preferably 300 ° C. or higher. More preferably, it is 450 ° C. or higher.

  In the present invention, in the finish rolling, lubrication rolling may be performed in order to reduce the hot rolling load. By performing lubrication rolling, there is an effect that the shape and material of the hot-rolled sheet are made more uniform. In addition, it is preferable to make the friction coefficient in the case of lubrication rolling into the range of 0.25-0.10. Moreover, the operation of hot rolling is further stabilized by combining lubrication rolling and continuous rolling.

The hot-rolled sheet that has been subjected to the above-described hot-rolling step is then subjected to pickling and cold-rolling in the cold-rolling step to become a cold-rolled plate.
The conditions for pickling may be generally known conditions and are not particularly limited. In addition, when the scale of a hot-rolled sheet is extremely thin, cold rolling may be performed immediately without performing pickling.
Further, the cold rolling conditions may be generally known conditions and are not particularly limited. Note that the cold rolling reduction is preferably 40% or more from the viewpoint of ensuring the uniformity of the structure.

Next, the cold-rolled sheet is subjected to a cold-rolled sheet annealing step by continuous annealing.
Further, following the cold-rolled sheet annealing step, temper rolling or leveler processing with an elongation of 1.0 to 15% may be performed. By performing temper rolling or leveler processing after the cold-rolled sheet annealing step, strain age hardening characteristics such as BH amount and ΔTS amount can be stably improved. The total elongation in temper rolling or leveler processing is preferably 1.0% or more. If the elongation is less than 1.0%, the strain age hardening property is not improved, while if the elongation exceeds 15%, the ductility of the steel sheet is lowered. Although the processing modes differ between temper rolling and leveler processing, the present inventors have confirmed that there is no significant difference in the effect on the strain age hardening characteristics of the steel sheet.

In the first preferred embodiment of the present invention, the slab composition includes Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al is 0.3 or more, C: 0.15% or less, Si: 2.0% In the following, Mn: 3.0% or less, P: 0.08% or less, S: 0.02% or less, or the following a group to d group a group: one or more of Cu, Ni, Cr, Mo in total 1.0% or less Group b: One or more of Nb, Ti, and V is 0.1% or less in total. Group c: B is 0.0030% or less. Group d: One or two of Ca and REM are combined in a total of 0.0010 to 0.010. %
Preferably, a slab composed of the remaining Fe and unavoidable impurities is used, and after the hot rolling step and the cold rolling step described above, a cold rolled plate is obtained, and then the cold rolled plate is reused. Annealing at a temperature not lower than the crystal temperature and not higher than 900 ° C. and holding time: 10 to 60 s, and then cooling to a temperature range of 500 ° C. or lower at a cooling rate of 10 to 300 ° C./s, and then the primary cooling. A cold-rolled sheet annealing step is performed in order to perform secondary cooling in which the residence time in the temperature range of 400 ° C. or lower is 300 s or lower.

The reason for limiting the slab composition in the first preferred embodiment of the present invention is the same as the reason for limiting the steel plate composition of the first preferred embodiment of the present invention. Moreover, in the 1st suitable aspect of this invention, it is more preferable that the coiling temperature in the said hot rolling process shall be 650 degrees C or less from a viewpoint of ensuring intensity | strength. Next, the reason for limiting the cold-rolled sheet annealing step in the first preferred embodiment of the present invention will be described.
Continuous annealing temperature: above the recrystallization temperature and below 900 ° C. The annealing temperature for continuous annealing is preferably set above the recrystallization temperature.

  If the continuous annealing temperature is lower than the recrystallization temperature, the recrystallization is not completed, and the strength satisfies the target, but the ductility is low, so that the formability is lowered and the steel sheet cannot be applied. In order to further improve the formability, the continuous annealing temperature is preferably 700 ° C. or higher. On the other hand, when the continuous annealing temperature exceeds 900 ° C., nitrides such as AlN are precipitated, and the amount of solute N in the steel plate as a product is insufficient. For this reason, the continuous annealing temperature is preferably not less than the recrystallization temperature and not more than 900 ° C.

Holding time at continuous annealing temperature: 10-60s
The holding time at the continuous annealing temperature is preferably as short as possible from the viewpoint of refining the structure and securing an amount of solute N that is higher than desired, but is preferably 10 s or longer from the viewpoint of operational stability. . If the holding time exceeds 60 s, it becomes difficult to refine the structure and secure the amount of solute N. For this reason, the holding time at the continuous annealing temperature is preferably in the range of 10 to 60 s.

Primary cooling: Cooling rate to temperature range below 500 ° C: 10-300 ° C / s
Cooling after soaking in continuous annealing is important from the viewpoint of refining the structure and securing the amount of dissolved N. In the present invention, cooling is performed at a temperature of 10 to 300 ° C./s as a primary cooling to a temperature range of 500 ° C. or lower. Cool continuously at speed. When the cooling rate is less than 10 ° C./s, it is difficult to ensure a uniform and fine structure and a desired amount of solid solution N or more. On the other hand, when the cooling rate exceeds 300 ° C./s, the uniformity of the material in the width direction of the steel sheet is insufficient. When the cooling stop temperature when cooling at a cooling rate of 10 to 300 ° C./s exceeds 500 ° C., the structure cannot be refined.

Secondary cooling condition: Cooling in which the residence time in the temperature range of 400 ° C. or higher of the primary cooling is 300 s or less is the secondary cooling after the primary cooling is important from the viewpoint of strain age hardening characteristics. Although the detailed mechanism is unknown at present, it is presumed that the amount of solid solution C and N varies depending on the secondary cooling conditions, and this affects the strain aging characteristics. In the present invention, following the primary cooling, it is preferable to continue the cooling, and to perform the cooling so that the residence time in the temperature range of 400 ° C. or higher below the primary cooling stop temperature is 300 s or shorter. In the present invention, after the continuous annealing, so-called overaging treatment may be performed, but when overaging treatment is performed, strain age hardening characteristics are deteriorated. Therefore, in the present invention, when passing the overaging zone of the continuous annealing furnace, it is desirable to set the temperature of the overaging zone to an extremely low temperature. In the first preferred embodiment of the present invention, the above-described temper rolling or leveler processing may be performed.

In the second preferred embodiment of the present invention, the slab composition contains Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al is 0.3 or more, C: 0.15% or less, Mn: 3.0% or less, S: 0.02% or less, and further containing one or two of Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0%, or further j group to m group j group : Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01%, one or more kinds k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: 0.01 to 0.2 Group 1: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 or 2 types m Group: 1 or 2 types of Ca and REM in total 0.0010 to 0.010%
Including one group or two or more groups selected from the above, preferably using a slab composed of the remaining Fe and inevitable impurities, and through the hot rolling step described above and the cold rolling step described above, After that, the cold-rolled sheet is annealed at a temperature of (Ac ₁ transformation point) to (Ac ₃ transformation point) and a holding time of 10 to 120 s, and then the average cooling rate between 600 to 300 ° C. is 1) or (2) When formula B <0.0003%
log CR = -1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.95 (1)
When B ≧ 0.0003%
log CR = −1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.40 …… (2)
(Here, CR: cooling rate (° C./s), Mn, Mo, Cr, Si, P, Cu, Ni: content of each element (mass%))
And a cold-rolled sheet annealing step for cooling at a critical cooling rate CR defined by (1) or higher.

The reason for limiting the slab composition in the second preferred embodiment of the present invention is the same as the reason for limiting the second preferred embodiment of the present invention.
The reason for limitation of the cold rolled sheet annealing step in the second preferred embodiment of the present invention will be described.
Annealing temperature: (Ac ₁ transformation point) to (Ac ₃ transformation point)
The annealing is preferably continuous annealing from the viewpoint of productivity. In the annealing treatment, heating is performed to a temperature in a two-phase region from (Ac ₁ transformation point) to (Ac ₃ transformation point). By heating to the two-phase region, the austenite (γ) phase and the ferrite (α) phase become two phases, C concentrates in the γ phase, and during cooling, the γ phase transforms into the martensite phase, This forms a composite structure of α + martensite. Thereby, ductility and workability are improved, and a low yield ratio is realized.

On the other hand, if the annealing temperature is less than the Ac ₁ transformation point, it becomes a ferrite + pearlite structure, and if it exceeds the Ac ₃ transformation point, the alloy element is not sufficiently concentrated in the γ phase, making it difficult to obtain a composite structure of α + martensite. Is insufficient.
Annealing holding time: 10 to 120 s
The holding time at the annealing temperature is preferably 10 to 120 s. The holding time at the annealing temperature is preferably as short as possible from the viewpoint of refining the structure and ensuring the amount of solute N, but is preferably 10 s or longer from the viewpoint of operational stability. When the holding time exceeds 120 s, it becomes difficult to refine the structure and secure the amount of solute N.

In addition, it is preferable that the heating to the soaking temperature of annealing shall be a heating rate of at least 5 ° C./s between 600 ° C. and (Ac ₁ transformation point). If it is less than 5 ° C./s, there is a problem in securing the amount of dissolved N. More preferably, it is 5-30 degreeC / s.
Cooling after soaking: The average cooling rate between 600 and 300 ° C is higher than the critical cooling rate CR. Cooling after soaking in annealing is important from the viewpoint of refining the structure, securing the amount of solute N and forming martensite. Yes, in the present invention, the average cooling rate between 600-300 ° C. is the following (1) or (2) when the formula B <0.0003% according to the amount of alloying elements
log CR = -1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.95 (1)
When B ≧ 0.0003%
log CR = −1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.40 …… (2)
(Here, CR: cooling rate (° C./s), Mn, Mo, Cr, Si, P, Cu, Ni: content of each element (mass%))
Cooling is performed at a critical cooling rate CR defined by In the formulas (1) and (2), elements not contained are calculated as 0.

  By cooling at an average cooling rate equal to or higher than the critical cooling rate CR in either of the formulas (1) or (2) according to the amount of alloy elements, precipitation of pearlite during cooling can be prevented. When cooling at a cooling rate less than CR (° C./s) defined by the above formulas, it becomes difficult to make the second phase martensite M (which may include partly bainite B). The tissue cannot be a composite tissue composed of α + M (+ B). When the average cooling rate exceeds 300 ° C./s, the material uniformity in the width direction of the steel sheet is insufficient. For this reason, the cooling after annealing is performed at an average cooling rate of 600 to 300 ° C. above CR defined by the formula (1) or (2), preferably 300 ° C./s or less. The average cooling rate in the temperature region below 300 ° C. is preferably 5 ° C./s or more.

  Also in the second preferred embodiment of the present invention, as in the first preferred embodiment, after the cold-rolled sheet annealing step, temper rolling or leveler processing with an elongation of 1.0 to 15% is further performed. Good.

(Example 1)
Molten steel having the composition shown in Table 1 was melted in a converter and made into a slab by a continuous casting method. These slabs were heated under the conditions shown in Table 2, roughly rolled to obtain sheet bars having the thicknesses shown in Table 2, and then hot-rolled sheets were formed by a hot rolling process in which finish rolling under the conditions shown in Table 2 was performed. In some cases, lubrication rolling was performed by finish rolling.

These hot-rolled sheets were made into cold-rolled sheets by a cold rolling process comprising pickling and cold rolling under the conditions shown in Table 2. Subsequently, these cold-rolled sheets were subjected to continuous annealing using a continuous annealing furnace under the conditions shown in Table 2. About some, temper rolling was performed following the cold-rolled sheet annealing process. In addition, all the annealing temperatures of continuous annealing were more than the recrystallization temperature.
The obtained cold-rolled annealed sheet was examined for the amount of dissolved N, microstructure, tensile characteristics, formability, strain age hardening characteristics, fatigue resistance characteristics, and impact resistance characteristics.
(1) Investigation of solute N amount The solute N amount was obtained by subtracting the precipitated N amount from the total N amount in steel obtained by chemical analysis. The amount of precipitated N was determined by an analysis method using the above-described constant potential electrolysis method.
(2) Microscopic structure A specimen is collected from each cold-rolled annealed plate, and a microscopic structure is imaged using an optical microscope or a scanning electron microscope with respect to a cross section (C cross section) orthogonal to the rolling direction. Was used to determine the structure fraction of the main phase ferrite, the type of the second phase, or the structure fraction.

Further, the crystal grain size of ferrite as the main phase is a value calculated by a quadrature method prescribed in ASTM from a structural photograph of a cross section (C cross section) orthogonal to the rolling direction or a nominal value obtained by a cutting method prescribed in ASTM. The larger of the particle sizes was adopted.
(3) Tensile properties JIS No. 5 test specimens were taken from each cold-rolled annealed sheet in the rolling direction, and subjected to a tensile test at a strain rate of 3 × 10 ⁻³ / s according to the provisions of JIS Z 2241 to yield. Strength YS, tensile strength TS, and elongation El were determined.
(4) Formability The r value was determined as an index of formability.

JIS 13B test pieces were collected from the rolling direction (L direction) of each cold-rolled annealed plate, 45 ° direction (D direction) with respect to the rolling direction, and 90 ° direction (C direction) with respect to the rolling direction. Obtain the width strain and plate thickness strain of each test piece when 15% uniaxial tensile pre-strain is applied to these test pieces, and the ratio of width strain to plate thickness strain,
r = ln (w / w ₀ ) / ln (t / t ₀ )
(Wherein, w _0, t ₀ is the width and thickness of the specimen before the test, w, t is the width and thickness of the test piece after the test.)
The r value in each direction is obtained from the following equation: r _mean = (r _L +2 r _D + r _c ) / 4
The average r value r _mean was determined by Here, r _L is the r value in the rolling direction (L direction), r _D is the r value in the rolling direction (L direction) with respect to the 45 ° direction (D direction), r _c is the rolling direction The r value is in the 90 ° direction (C direction) with respect to the (L direction).
(5) Strain age hardening characteristics JIS No. 5 test specimens were taken from each cold-rolled annealed sheet in the rolling direction, and as a pre-deformation, 5% tensile pre-strain was applied here, and then it was equivalent to 170 ° C x 20min. After performing the heat treatment, a tensile test was carried out at a strain rate of 3 × 10 ⁻³ / s, and the tensile properties (yield stress YS _BH , tensile strength TS _BH ) after pre-deformation-paint baking treatment were obtained, and BH Amount = YS _BH −YS _5% , ΔTS = TS _BH −TS was calculated. YS _5% is the deformation stress when the product plate is pre-deformed 5%, YS _BH and TS _BH are the pre-deformation-yield stress and tensile strength after paint baking, TS is the product plate Tensile strength.
(6) Fatigue resistance characteristics Fatigue specimens were taken from each cold-rolled annealed sheet in the rolling direction, and subjected to a tensile fatigue test with a minimum stress of 0 MPa in accordance with the provisions of JIS Z 2273. : 10 ⁷ times) σ _FL was determined. In addition, after applying a tensile pre-strain of 5% as pre-deformation, and then performing a heat treatment equivalent to 170 ° C x 20min, the same fatigue test was conducted to determine the fatigue limit (σ _FL ) _BH. The improvement margin ((σ _FL ) _BH −σ _FL ) of fatigue resistance by deformation-paint baking treatment was evaluated.
(7) Impact resistance properties Impact test specimens were taken from each cold-rolled annealed sheet in the rolling direction, and conformed to the high-speed tensile test method described in "Journal of the Society of Materials Science Japan, 10 (1998), p1058". Then, a high-speed tensile test was performed at a strain rate of 2 × 10 ³ / s, and a stress-strain curve was measured. Using the obtained stress-strain curve, the absorbed energy E was determined by integrating the stress in the range of strain 0-30%. Also, after applying a tensile pre-strain of 5% as a pre-deformation, and then performing a heat treatment equivalent to a 170 ° C x 20 min paint baking process, the same impact test was conducted to determine the absorbed energy E _BH and pre-deformation The improvement cost E _BH / E of impact resistance characteristics by paint baking treatment was evaluated.

In addition, the hot dip galvanization was given to some steel plate surfaces, and it was set as the plated steel plate, and various characteristics were evaluated similarly.
These results are shown in Table 3.

In the examples of the present invention, all have excellent ductility and excellent strain age hardening characteristics, exhibit a significantly high BH amount and ΔTS, and have a large allowance for improving fatigue resistance and impact resistance by strain aging treatment. .
In addition, the characteristics of the plated steel sheet in which the hot dip galvanized surfaces were applied to the surface of No. 11 and No. 13 steel sheets were almost the same as the characteristics before plating.
Further, regarding the steel plate No. 1 as an example of the present invention and the steel plate No. 5 as a comparative example, the aging conditions were variously changed and the strain age hardening characteristics were investigated. The results are shown in Table 4. The test method was the same except that only the aging temperature and aging time were changed.

  In the steel plate No. 1 as an example of the present invention, BH amount 115 MPa and ΔTS60 MPa were obtained by aging treatment of 170 ° C. × 20 min, which is a standard aging condition, but even under a wide range of aging treatment conditions as shown in Table 4. A BH amount of 80 MPa or more and ΔTS of 40 MPa or more could be satisfied. On the other hand, in the comparative example, even when the aging temperature was changed in the range of 100 to 300 ° C., the large BH amount and ΔTS as in the present invention example were not shown.

That is, the steel sheet of the present invention can secure a high BH amount and ΔTS even under a wide range of aging conditions.
(Example 2)
The steel having the composition shown in Table 5 was made into a slab in the same manner as in Example 1, and the slab was heated under the conditions shown in Table 6 and roughly rolled into a 25 mm-thick sheet bar. It was set as the hot rolled sheet by the hot rolling process which performs finish rolling of conditions. In addition, the sheet | seat bar | burr which flanks on the finishing rolling entrance side after rough rolling was joined by the melt-pressing method, and was continuously rolled. In addition, the temperature of the seat bar was adjusted by using an induction heating type seat bar edge heater and a sheet bar heater at the width end portion and the length direction end portion of the seat bar.

These hot-rolled sheets were made into 1.6 mm-thick cold-rolled sheets by a cold rolling process consisting of pickling and cold rolling under the conditions shown in Table 6. Subsequently, these cold-rolled sheets were subjected to continuous annealing using a continuous annealing furnace under the conditions shown in Table 6. In addition, all the annealing temperatures of continuous annealing were made into the recrystallization temperature or more.
The obtained cold-rolled annealed plate was examined in the same manner as in Example 1 for the amount of solute N, microstructure, tensile characteristics, formability, strain age hardening characteristics, fatigue resistance characteristics, and impact resistance characteristics.

  The results are shown in Table 7.

All of the examples of the present invention have excellent strain age hardening characteristics, exhibit stable and remarkably high BH amount and ΔTS, despite fluctuations in production conditions, and fatigue resistance and resistance to strain aging treatment. The cost for improving impact characteristics is also large. Moreover, in the example of this invention, the board thickness precision and shape of the product steel plate improved by implementing continuous rolling and the longitudinal direction and width direction temperature adjustment of a sheet bar.
(Example 3)
Molten steel having the composition shown in Table 8 was melted in a converter and made into a slab by a continuous casting method. These slabs were heated under the conditions shown in Table 9, roughly rolled to obtain sheet bars having the thicknesses shown in Table 9, and then hot-rolled sheets by a hot rolling process in which finish rolling under the conditions shown in Table 9 was performed. In some cases (steel plates No. 5-2 and No. 5-3), lubrication rolling was performed by finish rolling. Moreover, about one part, the sheet | seat bar which precedes and finishes by the finish rolling entrance side after rough rolling was joined by the melt-pressing method, and was continuously rolled. In addition, the temperature of the seat bar was adjusted by using an induction heating type seat bar edge heater and a sheet bar heater at the width end portion and the length direction end portion of the seat bar.

These hot-rolled sheets were made into cold-rolled sheets by a cold rolling process comprising pickling and cold rolling under the conditions shown in Table 9. Subsequently, these cold-rolled sheets were subjected to annealing (continuous annealing) using a continuous annealing furnace under the conditions shown in Table 9, and further subjected to a cold-rolled sheet annealing step of cooling under the conditions shown in Table 9 after annealing. About some, temper rolling was performed following the cold-rolled sheet annealing process.
The obtained cold-rolled annealed plate was examined in the same manner as in Example 1 for the amount of dissolved N, microstructure, tensile properties, formability, strain age hardening properties, and impact resistance properties.

  These results are shown in Table 10.

  In the examples of the present invention, all exhibit excellent ductility and a low yield ratio, and further have excellent strain age hardening characteristics, exhibit a significantly higher amount of BH and ΔTS, and allowance for improvement in impact resistance by strain aging treatment. Is also big.

Claims (14)

  A composition containing, in mass%, Al: 0.02% or less, N: 0.0050 to 0.0250%, N / Al being 0.3 or more, and solid solution N containing 0.0010% or more, and a ferrite phase having an average crystal grain size of 10 μm or less A high-tensile cold-rolled steel sheet having a tensile strength of 440 MPa or more and excellent strain age hardening characteristics, characterized by having a structure containing at least 50% in area ratio. The composition is in weight percent,
C: 0.15% or less, Si: 2.0% or less,
Mn: 3.0% or less, P: 0.08% or less,
S: 0.02% or less, Al: 0.02% or less,
N: 0.0050-0.0250%
And N / Al is 0.3 or more, N in a solid solution state is contained by 0.0010% or more, and the balance is composed of Fe and inevitable impurities. steel sheet. The high-tensile cold-rolled steel sheet according to claim 2, further comprising one group or two or more groups of the following groups a to d in mass% in addition to the composition.
Group a: One or more of Cu, Ni, Cr, and Mo total 1.0% or less in total Group b: One or two or more of Nb, Ti, and V total 0.1% or less in total c Group: B 0.0030% or less d group: One or two of Ca and REM in total 0.0010 to 0.010% The composition is in weight percent,
C: 0.15% or less, Mn: 3.0% or less,
S: 0.02% or less, Al: 0.02% or less,
N: 0.0050-0.0250%
And Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0%, or N / Al is 0.3 or more, and N in a solid solution state is 0.0010% or more, The balance is a composition composed of Fe and inevitable impurities, and the structure includes a ferrite phase having an average crystal grain size of 10 μm or less in an area ratio of 50% or more, and further includes a martensite phase in an area ratio of 3% or more. The high-tensile cold-rolled steel sheet according to claim 1. 5. The high-tensile cold-rolled steel sheet according to claim 4, further comprising one group or two or more groups among the following j groups to m groups in addition to the composition.
Group j: Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01% or two or more k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: One or more of 0.01 to 0.2% Group 1: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 or 2 types m group: 1 or 2 types of Ca and REM in total 0.0010 ~ 0.010%   The high-tensile cold-rolled steel sheet according to any one of claims 1 to 5, wherein the high-tensile cold-rolled steel sheet has a thickness of 3.2 mm or less.   A high-tensile cold-rolled steel sheet obtained by electroplating or hot-plating the high-tensile cold-rolled steel sheet according to any one of claims 1 to 6.   A steel slab having a composition containing Al: 0.02% or less, N: 0.0050-0.0250%, and N / Al being 0.3 or more in terms of mass% is heated to a slab heating temperature: 1000 ° C. or more and subjected to rough rolling. A sheet bar, a hot rolling process in which the sheet bar is subjected to finish rolling at a finish rolling exit temperature of 800 ° C. or more, and a winding temperature is 750 ° C. or less to obtain a hot rolled sheet; Cold rolling process in which pickling and cold rolling are performed to form a cold rolled sheet, and cold rolled sheet annealing is performed in which the cold rolled sheet is annealed by holding at a predetermined temperature for a predetermined time, and then cooled at a predetermined cooling rate. A method for producing a high-tensile cold-rolled steel sheet having a tensile strength of 440 MPa or more and excellent strain age hardening characteristics, characterized by sequentially performing the steps. The steel slab is in mass%,
C: 0.15% or less, Si: 2.0% or less,
Mn: 3.0% or less, P: 0.08% or less,
S: 0.02% or less, Al: 0.02% or less,
N: 0.0050-0.0250%
And N / Al is 0.3 or more, or a steel slab having a composition including one or more groups selected from the following groups a to d, and in the cold-rolled sheet annealing step, The annealing that is held at a predetermined temperature for a predetermined time is annealed at a recrystallization temperature of 900 ° C. or lower and a holding time of 10 to 60 seconds, and the cooling after annealing is performed at a cooling rate to a temperature range of 500 ° C. or lower. : Primary cooling that cools at 10 to 300 ° C./s, and then secondary cooling in which the residence time in the temperature range of 400 ° C. or higher is equal to or lower than the primary cooling stop temperature is 300 s or shorter. The manufacturing method of the high-tensile cold-rolled steel sheet according to claim 8.
Group a: One or more of Cu, Ni, Cr, and Mo total 1.0% or less in total Group b: One or two or more of Nb, Ti, and V total 0.1% or less in total c Group: B 0.0030% or less d group: One or two of Ca and REM in total 0.0010 to 0.010% The steel slab is in mass%,
C: 0.15% or less, Mn: 3.0% or less,
S: 0.02% or less, Al: 0.02% or less,
N: 0.0050-0.0250%
And Mo: 0.05 to 1.0%, Cr: 0.05 to 1.0%, or N / Al is 0.3 or more, or further among the following groups j to m The steel slab having a composition including one group or two or more groups selected from the above, and the annealing to be held for a predetermined time at a predetermined temperature in the cold-rolled sheet annealing step, (Ac ₁ transformation point) to (Ac ₃ transformation) The critical cooling rate defined by the following formula (1) or (2) is the average cooling rate between 600 and 300 ° C. The method for producing a high-tensile cold-rolled steel sheet according to claim 8, wherein the cooling is set to CR or higher.
Group j: Si: 0.05 to 1.5%, P: 0.03 to 0.15%, B: 0.0003 to 0.01% or two or more k group: Nb: 0.01 to 0.1%, Ti: 0.01 to 0.2%, V: One or more of 0.01 to 0.2% Group 1: Cu: 0.05 to 1.5%, Ni: 0.05 to 1.5%, 1 or 2 types m group: 1 or 2 types of Ca and REM in total 0.0010 ~ 0.010%
When B <0.0003%
log CR = -1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.95 (1)
When B ≧ 0.0003%
log CR = −1.73 [Mn + 2.67Mo + 1.3Cr + 0.26Si + 3.5P + 0.05Cu + 0.05Ni] + 3.40 …… (2)
Where CR: cooling rate (° C./s)
Mn, Mo, Cr, Si, P, Cu, Ni: Content of each element (mass%)   The high-tensile cold-rolled steel sheet according to any one of claims 8 to 10, wherein after the finish rolling, cooling is started within 0.5 s, the steel sheet is rapidly cooled at a cooling rate of 40 ° C / s or more, and the winding is performed. Manufacturing method.   The high-tensile cold-rolled steel sheet according to any one of claims 8 to 11, further comprising temper rolling or leveler processing with an elongation of 1.0 to 15% following the cold-rolled sheet annealing step. Method.   The method for producing a high-tensile cold-rolled steel sheet according to any one of claims 8 to 12, wherein sheet bars that follow each other are joined between the rough rolling and the finish rolling.   Between the rough rolling and the finish rolling, one or both of a sheet bar edge heater for heating the width end portion of the sheet bar and a sheet bar heater for heating the length end portion of the sheet bar are used. The method for producing a high-tensile cold-rolled steel sheet according to any one of claims 8 to 13.