JP6123958B1 - High strength steel plate and manufacturing method thereof - Google Patents

Abstract

A high strength steel sheet having a high yield ratio and small anisotropy in tensile properties is obtained. The steel structure has a specific component composition, and the steel structure is composed of ferrite in an area ratio of 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and residual austenite: 0 to 2%. The particle size is 15.0 μm or less, includes Ti carbide and / or V carbide, the average particle size of the Ti carbide and / or V carbide is 5 to 50 nm, and the total precipitation amount of Ti carbide and V carbide is the volume fraction The strength of the steel sheet is 0.005 to 0.050%.

Description

  The present invention relates to a high-strength steel plate applied to automobile parts and the like and a method for manufacturing the same.

  A high strength steel plate is preferably used as a material for automobile parts and the like from the viewpoint of reducing the weight of the component by reducing the thickness of the material. For example, skeletal parts and anti-collision parts are required to be difficult to be deformed at the time of collision, that is, to have a high yield ratio in order to ensure the safety of passengers. On the other hand, in order to stably press-form without generating cracks, a high-strength steel sheet having excellent bendability is desired. In response to such demands, various steel plates and manufacturing techniques thereof have been disclosed so far.

  Patent Document 1 discloses a high-strength steel sheet containing Nb and Ti in a total amount of 0.01% by mass or more and containing ferrite having a recrystallization rate of 80% or more as a main phase and a method for manufacturing the same.

  Patent Document 2 discloses a high-strength steel sheet excellent in impact resistance including 20 to 50% by area of non-recrystallized ferrite as a steel structure and a manufacturing method thereof.

  In Patent Document 3, one or more of V, Ti, and Nb is added, the main phase is ferrite or bainite, the amount of precipitation of iron carbide at the grain boundary is limited to a certain value, and the maximum of the iron carbide A hot-dip hot-dip steel sheet having excellent stretch flangeability and a particle diameter of 1 μm or less and a method for producing the same are disclosed.

Japanese Patent No. 4740099 Japanese Patent No. 4995109 JP-A-6-322479

  However, in the technique of Patent Document 1, neither the holding temperature of 650 ° C. after hot rolling nor the holding time of 500-400 ° C. in cooling after soaking in a continuous annealing furnace is controlled. Since it is considered that the average particle diameter of important Ti carbide and / or V carbide cannot be controlled, a high-strength steel sheet having a high yield ratio and excellent bendability cannot be obtained.

  In the technique of Patent Document 2, a large amount of Nb or Ti is added and non-recrystallized ferrite is contained as a steel structure in an area ratio of 20% or more, so that a high strength steel sheet having a high yield ratio and excellent bendability cannot be obtained.

  In the technique of Patent Document 3, neither the holding temperature of 650 ° C. after hot rolling nor the holding time of 500-400 ° C. in cooling after soaking in a continuous annealing furnace is controlled, which is important in the present invention. Since it is considered that the average particle diameter of Ti carbide and / or V carbide cannot be controlled, a high strength steel sheet having a high yield ratio and excellent bendability cannot be obtained.

  This invention is made | formed in view of such a situation, The objective is to obtain the high strength steel plate which has the bendability excellent in the high yield ratio.

  The inventors of the present invention conducted intensive research to solve the above problems. As a result, in the steel structure mainly composed of ferrite, the average crystal grain size of ferrite is refined below a certain level, and the volume fraction and particle size (average particle size) of Ti carbide and / or V carbide are appropriately controlled. In order to achieve this, it is necessary to control the coiling temperature after hot rolling and the residence time and soaking temperature in the specified temperature range at the time of temperature increase during annealing to an appropriate range. I found it effective.

  The present invention has been made based on the above findings, and the gist thereof is as follows.

  [1] Component composition is mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.10% or less, S: 0.020% or less, Al: 0.01 to 0.10%, N: 0.010% or less, Ti: 0.100% or less (including 0%), V: 0.100% or less (0% And a total of 0.005 to 0.100% of Ti and V, the balance being Fe and inevitable impurities, and the steel structure is ferrite in area ratio: 90% or more, the total of pearlite and cementite: 0 to 10%, the total of martensite and residual austenite: 0 to 2%, the ferrite has an average crystal grain size of 15.0 μm or less, includes Ti carbide and / or V carbide, Or the average particle size of V carbide is 5 to 50 nm, High strength steel sheet total precipitation amount of i carbides and V carbides is 0.005 to 0.050% by volume.

  [2] The component composition further includes, by mass%, Cr: 0.3% or less, Mo: 0.3% or less, B: 0.005% or less, Cu: 0.3% or less, Ni: 0.00. The high-strength steel sheet according to [1], containing any one or more of 3% or less and Sb: 0.3% or less.

  [3] The high-strength steel plate according to [1] or [2], which has a galvanized layer on the surface of the high-strength steel plate.

  [4] The high-strength steel plate according to [3], wherein the galvanized layer is a hot-dip galvanized layer.

  [5] The high-strength steel sheet according to [4], wherein the hot-dip galvanized layer is an alloyed hot-dip galvanized layer.

  [6] The high-strength steel plate according to [3], wherein the galvanized layer is an electrogalvanized layer.

  [7] A method for producing a high-strength steel sheet according to [1] or [2], wherein the steel is hot-rolled, and after the hot-rolling, a residence time in a temperature range of finish rolling temperature to 650 ° C. is 10 A hot rolling process in which the steel sheet is cooled under conditions of less than a second and wound at 500 to 700 ° C., and a cold rolling process in which the hot rolled steel sheet obtained in the hot rolling process is cold rolled at a rolling rate of 75% or less. Then, the cold-rolled steel sheet obtained in the cold rolling step is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at the time of temperature increase in a residence time: 15 seconds or more, and a soaking temperature after the residence: 760 -880 degreeC, soaking time: The manufacturing method of the high strength steel plate which has soaking | uniform-heating on the conditions for 120 second or less, and cooling on the conditions whose residence time of 400-500 degreeC is 100 seconds or less.

  [8] The method for producing a high-strength steel plate according to [7], including a plating step of plating the cold-rolled steel plate after the annealing step.

  [9] The method for manufacturing a high-strength steel sheet according to [8], wherein the plating treatment is a hot dip galvanizing treatment.

  [10] The method for producing a high-strength steel plate according to [9], including an alloying step of alloying the cold-rolled steel plate after the plating step.

  [11] The method for producing a high-strength steel sheet according to [8], wherein the plating treatment is an electrogalvanizing treatment.

  In the present invention, the production conditions such as the component composition, the winding condition after hot rolling, the residence time in the predetermined temperature range at the time of temperature rise of annealing, and the soaking temperature are appropriately controlled. By this control, the steel structure intended by the present invention can be obtained, and as a result, it is possible to stably produce a high-strength steel sheet having excellent bendability at a high yield ratio required for applications such as automobile parts. become. The high-strength steel sheet of the present invention makes it possible to further reduce the weight of an automobile, and the utility value of the present invention in the automobile and steel industry is extremely large.

  Hereinafter, embodiments of the present invention will be described. In addition, this invention is not limited to the following embodiment.

  First, the outline | summary of the high strength steel plate of this invention is demonstrated.

  The high-strength steel sheet of the present invention has a tensile strength of 330 MPa to less than 500 MPa, a yield ratio of 0.70 or more, and can be tightly bent by 180 ° in U-bending. Since the yield ratio is 0.70 or more, the high-strength steel sheet of the present invention has a high yield ratio. Moreover, since 180 degree | times contact bending is possible in U bending process, the high strength steel plate of this invention has the outstanding bendability.

  In the present invention, in particular, Ti: 0.100% or less (including 0%), V: 0.100% or less (including 0%), and Ti and V in total 0.005% to 0.100% It is particularly important to have the following component composition.

  By adjusting the component composition and manufacturing conditions, the steel structure is composed of essential ferrite and optional pearlite, etc., and the ferrite has an average crystal grain size of 15.0 μm or less, including Ti carbide and / or V carbide, By adjusting so that the average particle diameter of carbide and / or V carbide is 5 to 50 nm and the total amount of precipitation of Ti carbide and V carbide is 0.005 to 0.050% by volume, high yield A high-strength steel sheet having excellent bendability in the ratio can be obtained.

  In the present invention, Ti carbide and V carbide include Ti carbonitride, V carbonitride and Ti, V composite carbonitride. Ti and V composite carbonitrides may be regarded as Ti carbides or V carbides, and the average particle diameter and the total volume ratio may be considered.

  As described above, not only the component composition but also the production conditions are important in order for the average crystal grain size of ferrite and the average particle size and precipitation amount of carbides (Ti carbide and / or V carbide) to satisfy desired conditions. Specifically, in the cooling after hot rolling, the residence time in the temperature range of the finish rolling temperature to 650 ° C. is set to 10 seconds or less, and the winding temperature is set to 500 to 700 ° C. In the annealing heating, the residence time in the temperature range of 650 to 750 ° C. is set to 15 seconds or longer, and then the temperature is soaked at 760 to 880 ° C. for 120 seconds or shorter. During cooling after winding, Ti carbide and / or V carbide is uniformly and finely precipitated, and after cold rolling, ferrite is recrystallized at a relatively low temperature to produce fine ferrite, soaking. It is thought that the coarsening of ferrite grains and Ti carbide and / or V carbide can be suppressed.

  The yield strength and the tensile strength are obtained by taking a JIS No. 5 tensile test piece so that the tensile direction is perpendicular to the rolling direction and performing a tensile test in accordance with JIS Z 2241. The bendability is determined by an adhesion bending test described in JIS Z 2248.

  The high-strength steel sheet of the present invention completed based on the above knowledge becomes a steel sheet having excellent bendability at a high yield ratio required for materials such as automobile parts.

  Next, the reason for limiting the composition of the present invention, the reason for limiting the steel structure, and the reason for limiting the manufacturing conditions will be described.

(1) Component composition The high-strength steel sheet of the present invention is in mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.00. 10% or less, S: 0.020% or less, Al: 0.01 to 0.10%, N: 0.010% or less, Ti: 0.100% or less (including 0%), and V: 0.100 % Or less (including 0%), and Ti and V are contained in a total amount of 0.005 to 0.100% or less.

  Moreover, the high-strength steel sheet of the present invention may further include, as an optional component, in mass%, Cr: 0.3% or less, Mo: 0.3% or less, B: 0.005% or less, Cu: 0.3% Hereinafter, any one or more of Ni: 0.3% or less and Sb: 0.3% or less may be contained.

  The balance other than the above is Fe and inevitable impurities.

  In the following description of the component composition, “%” means “% by mass”.

C: 0.02% to less than 0.10% C is an element effective for increasing yield strength and tensile strength because it becomes Ti carbide or V carbide or increases pearlite or martensite. . If the C content is less than 0.02%, the total precipitation amount of carbides does not fall within the desired range, so that the intended tensile strength of the present invention cannot be obtained. When the C content is 0.10% or more, pearlite and martensite are excessively generated, so the yield ratio is lowered and the bendability is lowered. For this reason, C content shall be 0.02%-less than 0.10%. Preferably it is 0.02 to 0.06%.

Si: Less than 0.10% Si is generally effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. However, when Si is added, the increase in tensile strength is greater than the yield strength due to a significant improvement in work hardening ability, the yield ratio is lowered, and the surface properties are degraded. For this reason, Si content shall be less than 0.10%. In addition, although the lower limit of Si content is not specifically limited, Since yield strength and tensile strength are raised also in structures other than Si, in this invention, it is so preferable that Si content is small. Therefore, in the present invention, Si may not be added, but Si may inevitably be contained in an amount of 0.005% in production.

Mn: less than 1.0% Mn is effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. However, if the Mn content is 1.0% or more, the martensite fraction in the steel structure increases, so the tensile strength increases excessively, and the intended tensile strength of the present invention cannot be obtained, yielding. Ratio and bendability are reduced. Therefore, the Mn content is less than 1.0%. Mn may not be added, but when Mn is added, the preferable Mn content for the lower limit is 0.2% or more. A preferable Mn content for the upper limit is 0.8% or less.

P: 0.10% or less P is effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. For this reason, in this invention, P can be contained suitably. However, if the P content exceeds 0.10%, the ferrite grain boundary becomes brittle due to casting segregation or ferrite grain boundary segregation, and the bendability decreases. Therefore, the P content is 0.10% or less. P may not be added, but when P is added, the preferable P content for the lower limit is 0.01% or more. A preferable P content for the upper limit is 0.04% or less.

S: 0.020% or less S is an element inevitably included as an impurity. Since the bendability and local elongation are reduced by the formation of inclusions such as MnS, it is preferable to reduce the S content as much as possible. In the present invention, the S content is 0.020% or less. Preferably, the content is 0.015% or less. In addition, as above-mentioned, S content is so preferable that it is low, and it is not necessary to add S in this invention. However, there are cases in which S is contained in an amount of 0.0003%.

Al: 0.01-0.10%
Al is added for deoxidation in the refining process and for fixing the solid solution N as AlN. In order to obtain a sufficient effect, the Al content needs to be 0.01% or more. On the other hand, if the Al content exceeds 0.10%, a large amount of AlN precipitates and the bendability decreases. Therefore, the Al content is set to 0.01 to 0.10%. Preferably, the content is 0.01 to 0.07%. More preferably, the content is 0.01 to 0.06%.

N: 0.010% or less N is an element inevitably mixed up to the hot metal refining process. If the N content exceeds 0.010%, Ti carbides and V carbides will precipitate during casting, and Ti carbides and V carbides will not dissolve and remain as coarse carbides during slab heating. Invite. Therefore, the N content is 0.010% or less. In the present invention, N may not be added, but may contain 0.0005% N in production.

Ti: 0.100% or less (including 0%)
V: 0.100% or less (including 0%)
Ti and V in total 0.005 to 0.100%
Ti and V are important elements that contribute to the refinement of the average ferrite grains and the increase in yield ratio due to the precipitation of Ti carbide and / or V carbide. If the total of Ti and V is less than 0.005%, the volume ratio of Ti carbide and / or V carbide becomes insufficient, and as a result, the precipitation amount of carbide is not in a desired range, and the effect of the present invention cannot be obtained. . On the other hand, if the total of Ti and V exceeds 0.100%, Ti carbide and / or V carbide precipitates excessively, and unrecrystallized ferrite having poor ductility remains even after annealing, so that the bendability deteriorates. Therefore, Ti and V are Ti: 0.100% or less (including 0%) and V: 0.100% or less (including 0%), and Ti and V are 0.005 to 0.100% in total. To do. A preferable total amount for the lower limit is 0.007% or more, and a preferable total amount for the upper limit is 0.040%.

  The high-strength steel sheet of the present invention can contain the following components as optional components.

Cr: 0.3% or less Cr may be contained as a trace element that does not impair the effects of the present invention. If the Cr content exceeds 0.3%, the martensite may be excessively generated due to the improvement of the hardenability, leading to a decrease in the yield ratio. Therefore, when adding Cr, Cr content shall be 0.3% or less.

Mo: 0.3% or less Mo may be contained as a trace element that does not impair the effects of the present invention. However, if the Mo content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, which may lead to a decrease in yield ratio. Therefore, when Mo is added, the Mo content is 0.3% or less.

B: 0.005% or less B may be contained as a trace element that does not impair the effects of the present invention. However, if the B content exceeds 0.005%, martensite may be generated excessively due to the improvement of hardenability, leading to a decrease in yield ratio. Therefore, when adding B, B content shall be 0.005% or less.

Cu: 0.3% or less Cu may be contained as a trace element that does not impair the effects of the present invention. However, if the Cu content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, which may lead to a decrease in yield ratio. Therefore, when adding Cu, Cu content shall be 0.3% or less.

Ni: 0.3% or less Ni may be contained as a trace element that does not impair the effects of the present invention. However, if the Ni content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, leading to a decrease in yield ratio. Therefore, when Ni is added, the Ni content is 0.3% or less.

Sb: 0.3% or less Sb may be contained as a trace element that does not impair the effects of the present invention. However, if the Sb content exceeds 0.3%, the high-strength steel sheet is brittle and the bendability deteriorates. Therefore, when adding Sb, Sb content shall be 0.3% or less.

  The balance other than the above is Fe and inevitable impurities. Further, in the present invention, in addition to the above, elements such as Nb, Sn, Co, W, Ca, Na, and Mg may be contained as inevitable impurities within a minute range that does not impair the effects of the present invention. . The “trace range” means 0.01% or less of these elements in total.

(2) Steel structure The steel structure of the high-strength steel sheet of the present invention is composed of ferrite by area ratio of 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and residual austenite: 0 to 3%. . In this steel structure, the average crystal grain size of the ferrite is 15.0 μm or less, the average particle size of Ti carbide and / or V carbide is 5 to 50 nm, and the precipitation amount of Ti carbide and / or V carbide Is a volume ratio of 0.005 to 0.050%.

Ferrite: 90% or more Ferrite has good ductility and is contained in the steel structure as a main phase, and its content is 90% or more in terms of area ratio. If the ferrite content is less than 90% by area ratio, the high yield ratio intended by the present invention cannot be obtained, and the anisotropy of tensile properties also increases. Therefore, the ferrite content is 90% or more in terms of area ratio. Preferably, it is 95% or more. The steel structure of the high-strength steel sheet of the present invention may be a ferrite single phase (the ferrite content is 100% in area ratio).

Total of pearlite and cementite: 0-10%
Pearlite and cementite are effective in obtaining desired yield strength and tensile strength. However, if the total of pearlite and cementite exceeds 10% in terms of area ratio, the high yield ratio intended by the present invention cannot be obtained, and the anisotropy of tensile properties increases. Therefore, the total of pearlite and cementite is 0 to 10% in terms of area ratio. Preferably it is 0 to 5%.

Total of martensite and retained austenite: 0-3%
The steel structure may contain 0 to 3% of martensite and retained austenite in terms of area ratio. When the sum of martensite and retained austenite exceeds 2%, a yield ratio of 0.70 or more cannot be obtained. Therefore, the total of martensite and retained austenite is 0 to 3%.

The average crystal grain size of ferrite is 15.0 μm or less It is important to adjust the average crystal grain size of ferrite to a desired range in order to obtain a high yield ratio of 0.70 or more, which is an object of the present invention. When the average crystal grain size of ferrite exceeds 15.0 μm, a yield ratio of 0.70 or more cannot be obtained. Therefore, the average crystal grain size of ferrite is 15.0 μm or less. Preferably it is 10.0 μm or less. The lower limit of the average ferrite grain size is not particularly limited, but if it is less than 1.0 μm, the tensile strength and yield strength increase excessively, leading to deterioration of bendability and elongation, so the average ferrite grain size is 1.0 μm. The above is preferable.

The average particle diameter of Ti carbide and / or V carbide is 5 to 50 nm.
Ti carbides and V carbides are mainly precipitated in ferrite grains, and the average particle size is important for achieving both the high yield ratio and excellent bendability that are the object of the present invention. If the particle diameter is less than 5 nm, not only the yield strength and the tensile strength are excessively increased, but also the bendability is lowered. If the particle diameter exceeds 50 nm, the increase in yield strength is insufficient, and the high yield ratio intended by the present invention cannot be obtained. Therefore, the average particle diameter of Ti carbide and / or V carbide is set to 5 to 50 nm. A preferable average particle size for the lower limit is 10 nm or more. A preferable average particle size for the upper limit is 40 nm or less. In the present invention, the average particle diameter is measured without distinguishing between Ti carbide and V carbide.

The total precipitation amount of Ti carbide and V carbide is 0.005 to 0.050% by volume.
It is important to adjust the precipitation amount of Ti carbide and V carbide within a desired range in order to achieve both the high yield ratio and excellent bendability that are the object of the present invention. If the total precipitation amount of Ti carbide and V carbide is less than 0.005% by volume, the increase in yield strength is insufficient, and the high yield ratio intended by the present invention cannot be obtained. When the total precipitation amount of Ti carbide and V carbide exceeds 0.050% by volume, recrystallization of ferrite is remarkably suppressed, yield strength and tensile strength increase excessively, and bendability decreases. . Moreover, when the said precipitation amount exceeds 0.050%, tensile strength will increase too much and it may remove | deviate from the range which this invention aims. Therefore, the total precipitation amount of Ti carbide and V carbide is 0.005 to 0.050% by volume. A preferable volume ratio for the lower limit is 0.010% or more. A preferable volume ratio for the upper limit is 0.040% or less. When Ti carbide is not included, Ti carbide is considered as 0, and when V carbide is not included, V carbide is considered as 0.

  In addition, the area ratio of each structure was observed by SEM in the range of the plate thickness 1/8 to 3/8 centered on the 1/4 position in the plate thickness direction from the steel plate surface side of the cross section perpendicular to the rolling width direction. It is determined by the point counting method described in E 562-05. For the average crystal grain size of ferrite, the range of the plate thickness 1/8 to 3/8 centered on the above-mentioned plate thickness 1/4 position is observed with SEM, and the equivalent circle diameter is calculated from the observation area and the number of crystal grains. Ask for. The average particle diameter of Ti carbide or V carbide is determined by preparing a thin film sample from a high-strength steel plate and calculating the equivalent circle diameter from the TEM observation image (calculating from the observation area and the number of particles). The total volume ratio of Ti carbide and V carbide is determined by the extraction residue method.

(3) Manufacturing conditions The high-strength steel sheet of the present invention melts steel having the above component composition, manufactures a slab (steel piece) by casting, performs hot rolling and cold rolling, and then anneals in a continuous annealing furnace. It is manufactured by. Pickling may be performed after hot rolling. Hereinafter, the manufacturing method of this invention which has a hot rolling process, a cold rolling process, and an annealing process is demonstrated. In the following description, temperature means surface temperature.

  The casting method is not particularly limited, and casting may be performed by either the ingot casting method or the continuous casting method as long as the segregation of the significant component composition and the unevenness of the structure do not occur.

  In hot rolling, a high-temperature cast slab may be rolled as it is, or the slab cooled to room temperature may be reheated and then rolled. If there is a surface defect such as a crack at the time of the slab, the slab can be treated with a grinder. When the slab is reheated, it is preferably heated to 1100 ° C. or higher in order to dissolve Ti carbide and / or V carbide.

  In the hot rolling process, steel is hot rolled, and after the hot rolling, the steel sheet is cooled under a condition that the residence time in the temperature range of finish rolling temperature to 650 ° C. is 10 seconds or less, and wound at 500 to 700 ° C. It is a process to take.

  In hot rolling, the slab is subjected to rough rolling and finish rolling. Then, the hot-rolled steel sheet is taken up as a hot rolled coil. The rough rolling conditions and finish rolling conditions in the hot rolling are not particularly limited, and may be determined according to a conventional method. When the finish rolling temperature is less than the Ar3 point, coarse ferrite stretched in the rolling direction is generated in the steel structure of the hot-rolled steel sheet, and ductility may be lowered after annealing. For this reason, it is preferable that finishing rolling temperature shall be Ar3 point or more. In addition, Ar3 point can be calculated | required by measuring the temperature which a ferrite transformation starts when it cools continuously at 1 degree-C / s from an austenite single phase temperature range using a transformation point measuring apparatus (for example, for master test machine). .

Retention time in the temperature range from the finish rolling temperature to 650 ° C .: 10 seconds or less In cooling after hot rolling, the average grain size of ferrite is controlled by appropriately controlling the residence time in the temperature range from the finish rolling temperature to 650 ° C. Can be suppressed. For this reason, the said cooling conditions are important in this invention. If the residence time in the temperature range of the finish rolling temperature to 650 ° C. exceeds 10 seconds in the cooling after finish rolling, coarse Ti carbides and V carbides excessively precipitate after winding of the hot rolling. Tends to become coarser, and the average crystal grain size of ferrite exceeds 15.0 μm, so the yield ratio decreases. Therefore, the residence time in the temperature range of the finish rolling temperature to 650 ° C. in the cooling is 10 seconds or less. In addition, although the minimum of the said residence time is not specifically limited, It is preferable to retain for 1 second or more from a viewpoint which precipitates Ti carbide | carbonized_material and V carbide | carbonized_material uniformly at the time of annealing, and makes a ferrite crystal grain diameter uniform. In addition, the lower limit of the temperature range in which the residence time is controlled is 650 ° C. because the average particle diameter of Ti carbide or the like is out of the scope of the present invention, or the total amount of precipitation of Ti carbide is out of the scope of the present invention. To do.

Winding temperature: 500-700 ° C
The coiling temperature is important in order to control the ferrite average crystal grain size after annealing to 15.0 μm or less by adjusting the precipitation amount of Ti carbide and V carbide and the average particle diameter thereof. When the coiling temperature is less than 500 ° C at the center in the width direction of the steel sheet, the carbides do not sufficiently precipitate during cooling after winding, coarse carbides precipitate during annealing and soaking, and the ferrite grain size increases. Therefore, a high yield ratio cannot be obtained, and the tensile strength is also reduced. When the coiling temperature exceeds 700 ° C., coarse Ti carbides and V carbides precipitate during cooling after winding, and the ferrite grain size becomes coarse during annealing, so a high yield ratio cannot be obtained, and the tensile strength is also high. Get smaller. Therefore, the coiling temperature is 500 to 700 ° C. A preferable coiling temperature for the lower limit is 550 ° C. or higher. A preferable coiling temperature for the upper limit is 650 ° C. or less.

  The cold rolling step is a step of cold rolling the hot rolled steel sheet obtained in the hot rolling step. The rolling rate of cold rolling is 75% or less. Preferably it is 30 to 75%. If the rolling rate exceeds 75%, the average particle diameter of the carbide becomes coarse and the desired bendability cannot be obtained, so 75% or less is necessary. A rolling rate of 30% or more is preferable because excellent bendability can be obtained by completely recrystallizing ferrite during annealing.

  Annealing consists of a process of raising the temperature to a soaking temperature and then cooling it using a continuous annealing furnace. The annealing process in the present invention means that the cold-rolled steel sheet obtained in the cold rolling process is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at a temperature rise time of 15 seconds or more, and after the retention Soaking temperature: 760 to 880 ° C., soaking time; soaking under conditions of 120 seconds or less, and after the soaking, cooling is performed under conditions of a residence time in a temperature range of 400 to 500 ° C. of 100 seconds or less.

Residence time in the temperature range of 650 to 750 ° C. at the time of temperature rise: 15 seconds or more The residence time at 650 to 750 ° C. at the time of temperature rise is important for controlling the ferrite average crystal grain size after annealing to 15.0 μm or less. Manufacturing conditions. If the residence time at 650 to 750 ° C. at the time of temperature rise is less than 15 seconds, recrystallization of the ferrite will not be completed during the temperature rise, so recrystallization proceeds during relatively high temperature soaking and the average ferrite grain size is coarse Turn into. Therefore, the residence time at 650 to 750 ° C. when the temperature is raised is 15 seconds or more. Preferably, the residence time at 650 to 750 ° C. during the temperature rise is 20 seconds or more. The upper limit of the residence time is not particularly limited, but if the residence time is too long, the Ti carbide and V carbide are coarsened, so the residence time is preferably 300 seconds or less.

Soaking temperature: 760 to 880 ° C., soaking time: 120 seconds or less Soaking temperature and soaking time are important conditions for controlling the ferrite average crystal grain size. If the soaking temperature is less than 760 ° C., the recrystallization of ferrite becomes insufficient and the bendability deteriorates. When the soaking temperature exceeds 880 ° C., the ferrite average crystal grain size becomes coarse, the desired yield ratio of the present invention cannot be obtained, and the tensile strength is also reduced. Therefore, the soaking temperature is set to 760 to 880 ° C. On the other hand, if the soaking time exceeds 120 seconds, the average grain size of ferrite becomes coarse, so that the intended tensile strength and high yield ratio of the present invention cannot be obtained. For this reason, the soaking time is 120 seconds or less. Preferably it is 60 seconds or less. The lower limit of the soaking time is not particularly limited, but it is preferable that the soaking time is 30 seconds or longer because it is preferable to completely recrystallize ferrite from the viewpoint of bendability.

  The heating method at the time of temperature increase and soaking is not particularly limited, and can be performed by a radiant tube method, a direct fire heating method, or the like.

  The cooling condition in the cooling after soaking is that the residence time in the temperature range of 400 to 500 ° C. is 100 seconds or less. It is necessary for the residence time to be 100 seconds or less to make the average particle size of the carbides 50 nm or less. The lower limit of the residence time is not particularly limited. However, if it is extremely short, the solid solution C in the ferrite is increased and the aging resistance is deteriorated or excessive investment in the cooling equipment is required. . More preferably, it is 10 seconds or more. Here, “the residence time in the temperature range of 400 to 500 ° C.” means the total time during which the steel plate being cooled is at a temperature of 400 to 500 ° C. If the cooling stop temperature is 400 ° C. or higher. Means the total time from the cooling stop temperature to 500 ° C. Moreover, the residence in this temperature range is equivalent to an overaging treatment. In addition, although other cooling conditions are not specifically limited, Cooling stop temperature is 400-500 degreeC, and the conditions whose average cooling rate is 30 degrees C / s or less are mentioned.

  Plating can be applied to the surface of the high-strength steel plate obtained as described above. The plating is preferably galvanization, and the galvanization layer is formed on the high-strength steel plate by applying galvanization to the high-strength steel plate of the present invention. Among galvanizing (electrogalvanizing, hot dip galvanizing, etc.), hot dip galvanizing immersed in a hot dip galvanizing bath is suitable.

  An alloyed hot-dip galvanized layer is formed by subjecting a hot-dip galvanized layer formed by hot-dip galvanizing to a high-strength steel plate. When the alloying treatment is performed, if the holding temperature is less than 450 ° C., alloying does not proceed sufficiently, and the plating adhesion and corrosion resistance may deteriorate. Further, when the holding temperature exceeds 560 ° C., alloying proceeds excessively, and problems such as powdering may occur during pressing. Therefore, the holding temperature is preferably 450 to 560 ° C. Further, if the holding time is less than 5 seconds, alloying does not proceed sufficiently and the plating adhesion and corrosion resistance may deteriorate, so the holding time is preferably 5 seconds or more.

  Thereafter, temper rolling with an elongation of 0.1 to 5.0% may be performed as necessary.

  As described above, the high-strength steel plate targeted by the present invention is obtained. Even if the high-strength steel sheet of the present invention is subjected to a surface treatment such as a chemical conversion treatment or an organic coating treatment or a coating, the target properties of the present invention are not impaired.

  Hereinafter, the present invention will be described in detail with reference to examples.

  After steel slabs of steels A to O having the composition shown in Table 1 were soaked at 1250 ° C. for 1 hour, after rolling at a finish plate thickness of 3.2 mm and a finish rolling temperature of 900 ° C. of Ar 3 or higher, Table 2 The sample was cooled under the conditions shown in FIG. The picked hot-rolled steel sheet was pickled and then cold-rolled to a finished sheet thickness of 1.4 mm to obtain a cold-rolled steel sheet, and annealed under the conditions shown in Table 2 to obtain No. 1 to 31 high strength steel sheets were produced. In addition, the cooling conditions for cooling in annealing are as follows: the cooling stop temperature is 480 ° C., the average cooling rate is 20 ° C./s or less, and the residence time in the temperature range of 400 to 500 ° C. (500 ° C. to the temperature range of cooling stop temperature) is 30 seconds. It was. Annealing was performed using CAL when plating was not performed. Moreover, when performing plating, CGL was used, and hot dip galvanization or galvannealing was performed. When the plated layer was an alloyed hot-dip galvanized layer, an alloying treatment was performed by holding at 510 ° C. for 10 seconds.

  Steel structure observation and a tensile test were performed on the obtained high-strength steel sheet.

  As for the area ratio of each steel structure, the area ratio of each structure is an SEM in the range of 1/8 to 3/8 of the sheet thickness centered at 1/4 position in the sheet thickness direction from the steel sheet surface side of the cross section perpendicular to the rolling width direction. And determined by the point counting method described in ASTM E 562-05. For the average crystal grain size of ferrite, the range of the plate thickness 1/8 to 3/8 centered on the above-mentioned plate thickness 1/4 position is observed with SEM, and the equivalent circle diameter is calculated from the observation area and the number of crystal grains. I asked for it. The average particle diameter of the carbides (Ti carbide, V carbide) was observed by TEM, and the equivalent circle diameter was determined by image processing. The total volume ratio of Ti carbide and V carbide was determined by the extraction residue method. All observations were made with 10 fields of view, and the average was calculated. The results are shown in Table 2. In Table 2, α is ferrite, P is pearlite, M is martensite, θ is cementite, α grain size is the average ferrite grain size, and M (C, N ) Particle size means the average particle size of carbides, and M (C, N) volume fraction means the total amount of precipitation of Ti carbides and V carbides. Moreover, M in the above M (C, N) means Ti or V.

  The tensile strength (TS) and the yield ratio (YR) were determined by a tensile test in accordance with JIS Z 2241 using a JIS No. 5 tensile test specimen collected so that the tensile direction was perpendicular to the rolling direction. Further, the bending test was performed in accordance with JIS Z 2248 by collecting a test piece in a direction in which a bending ridge line is parallel to the rolling direction. The tensile strength of 330 MPa to less than 500 MPa, the yield ratio of 0.70 or more, and contact bending were performed at N = 3, and it was evaluated that no cracks occurred.

Table 2 shows the observation results of the steel structure, the tensile test results, and the bending test results.
No. 1 to 3, 6, 8, 9, 14 to 16, 18, 19, 22, 24, 25, and 28 satisfy all the requirements of the present invention. An excellent high-strength steel sheet is obtained.
On the other hand, no. 4, 5, 7, 10-13, 17, 20, 21, 23, 26, 27, 29, 30-31 have component compositions or production conditions outside the scope of the present invention, and a desired steel structure is obtained. Therefore, the high strength steel plate intended by the present invention is not obtained.

The high-strength steel sheet of the present invention is suitable for a field that requires high yield ratio and isotropy of tensile properties, mainly for automobile inner plate parts.

Claims (11)

Component composition is mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.10% or less, S: 0.020 %: Al: 0.01-0.10%, N: 0.010% or less, Ti: 0.100% or less (including 0%), V: 0.100% or less (including 0%) and It contains 0.005 to 0.100% in total of Ti and V, and the balance consists of Fe and inevitable impurities,
Steel structure consists of ferrite in area ratio: 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and retained austenite: 0 to 3%,
The average crystal grain size of the ferrite is 15.0 μm or less,
Ti carbide and / or V carbide is included, and the average particle diameter of the Ti carbide and / or V carbide is 5 to 50 nm,
A high-strength steel sheet in which the total precipitation amount of Ti carbide and V carbide is 0.005 to 0.050% by volume.   The component composition is further in mass%, Cr: 0.3% or less, Mo: 0.3% or less, B: 0.005% or less, Cu: 0.3% or less, Ni: 0.3% or less Sb: The high-strength steel plate according to claim 1 containing one or more of 0.3% or less.   The high-strength steel sheet according to claim 1 or 2, wherein the surface has a galvanized layer.   The high-strength steel sheet according to claim 3, wherein the galvanized layer is a hot-dip galvanized layer.   The high-strength steel sheet according to claim 4, wherein the hot-dip galvanized layer is an alloyed hot-dip galvanized layer.   The high-strength steel sheet according to claim 3, wherein the galvanized layer is an electrogalvanized layer. It is a manufacturing method of the high strength steel plate according to claim 1 or 2,
Hot rolling the steel, and after the hot rolling, a hot rolling step of cooling the steel plate under a condition that the residence time in the temperature range of the finish rolling temperature to 650 ° C. is 10 seconds or less and winding at 500 to 700 ° C .;
A cold rolling step of cold rolling the hot rolled steel sheet obtained in the hot rolling step at a rolling rate of 75% or less;
The cold-rolled steel sheet obtained in the cold rolling step is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at the time of temperature increase for a residence time of 15 seconds or more, and after the residence, a soaking temperature: 760 to 880. C., soaking time: An annealing step of soaking under conditions of 120 seconds or less and cooling under conditions where the residence time in the temperature range of 400 to 500.degree. C. is 100 seconds or less.   The manufacturing method of the high strength steel plate of Claim 7 which has a plating process of plating the cold-rolled steel plate after the said annealing process.   The method for producing a high-strength steel sheet according to claim 8, wherein the plating process is a hot dip galvanizing process.   The manufacturing method of the high strength steel plate of Claim 9 which has an alloying process of alloying the cold-rolled steel plate after the said plating process. The method for producing a high-strength steel sheet according to claim 8, wherein the plating process is an electrogalvanizing process.