JP5685166B2 - High-strength steel sheet with excellent workability and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength steel sheet excellent in workability having a tensile strength of 590 MPa or more and a method for producing the same.

In the automobile industry, there is an urgent need to deal with global environmental problems such as CO ₂ emission regulations. On the other hand, from the viewpoint of ensuring the safety of passengers, the collision safety standards for automobiles have been strengthened, and structural designs that can sufficiently ensure the safety in the riding space are being advanced. In order to achieve these requirements at the same time, it is effective to use a high-strength steel plate having a tensile strength of 590 MPa or more as a structural member of an automobile and further reduce the thickness thereof to reduce the weight of the vehicle body. However, generally, when the strength of the steel plate is increased, the workability deteriorates. Therefore, in order to apply the high strength steel plate to an automobile member, improvement of the workability is an unavoidable issue.

  As steel sheets having both strength and workability, DP (Dual Phase) steel sheets whose structure is composed of ferrite and martensite, and TRIP (Transformation Induced Plasticity) using transformation-induced plasticity of retained austenite (residual γ): ) Steel plates are known.

  Among these, as a TRIP steel plate having strength and workability, for example, a steel plate of Patent Document 1 is known. This document discloses a technique for improving the strength and workability (particularly, elongation) of a steel sheet by making the metal structure of the steel sheet a composite structure in which martensite and residual γ are mixed in ferrite.

  Patent Document 2 discloses a technique for improving the press formability by improving the balance of strength (TS) and elongation (EL) (specifically, TS × EL) for the TRIP steel sheet. In this document, in order to improve the press formability, the metal structure is a structure containing ferrite, residual γ, bainite and / or martensite. This document describes that the residual γ has an effect of improving the elongation of the steel sheet.

  As disclosed in Patent Documents 1 and 2 described above, by making the metal structure of the steel sheet a structure containing residual γ, the strength of the steel sheet can be increased and the elongation characteristics can be improved.

Japanese Patent Application Laid-Open No. 11-296991 JP 2007-126747 A

  Recently, the required properties for workability of steel sheets are becoming more and more severe. For example, steel sheets used for pillars and members are required to be stretched or drawn under severer conditions than ever before. Therefore, TRIP steel sheets are required to improve local deformability such as stretch flangeability (λ) and bendability (R) without deteriorating strength and elongation. However, the above-described TRIP steel sheet has a problem in that the residual γ is transformed into very hard martensite during processing, and therefore the local deformability such as stretch flangeability and bendability is inferior.

  The present invention has been made paying attention to the circumstances as described above, and the purpose thereof is a high strength steel sheet having a tensile strength of 590 MPa or more and excellent workability with improved both elongation and local deformability. An object of the present invention is to provide a high-strength steel sheet and a method for producing the same.

  The high-strength steel sheet according to the present invention that has solved the above problems is mass%, C: 0.10 to 0.3%, Si: 1.0 to 3%, Mn: 1.0 to 2. 5%, Al: 0.005 to 3%, P: 0.1% or less, S: 0.05% or less, with the balance being iron and inevitable impurities. The metal structure of the steel sheet includes polygonal ferrite, bainite, tempered martensite, and retained austenite. (1) When the metal structure is observed with a scanning electron microscope, (1a) the area ratio of the polygonal ferrite a is more than 50% with respect to the entire metal structure, and (1b) the bainite is a high-temperature region-generated bainite having an average interval of adjacent residual austenite and / or carbide of 1 μm or more, and adjacent residual austenite and / or It is comprised with the composite structure with the low temperature area | region production bainite whose average space | interval of a carbide | carbonized_material is less than 1 micrometer, The area ratio b of the said high temperature area production | generation bainite is 5 to 40% with respect to the whole metal structure, The total area ratio c with the tempered martensite satisfies 5 to 40% with respect to the entire metal structure, (2) The gist is that the volume fraction of the retained austenite measured by the saturation magnetization method is 5% or more with respect to the entire metal structure.

  When the steel structure has an MA mixed phase in which quenched martensite and retained austenite are present when the metal structure is observed with an optical microscope, the cross-section of the steel sheet is compared with the total number of MA mixed phases. The number ratio of MA mixed phases satisfying the equivalent circle diameter d of more than 7 μm is preferably less than 15% (including 0%).

  The average equivalent circle diameter D of the polygonal ferrite grains is preferably 10 μm or less (not including 0 μm).

The steel sheet, as another element,
(A) Cr: 1% or less (not including 0%) and / or Mo: 1% or less (not including 0%),
(B) A group consisting of Ti: 0.15% or less (not including 0%), Nb: 0.15% or less (not including 0%), and V: 0.15% or less (not including 0%) One or more elements selected from
(C) Cu: 1% or less (not including 0%) and / or Ni: 1% or less (not including 0%),
(D) B: 0.005% or less (excluding 0%),
(E) Ca: 0.01% or less (not including 0%), Mg: 0.01% or less (not including 0%), and rare earth elements: 0.01% or less (not including 0%) One or more elements selected from the group,
Etc. may be contained.

  The present invention includes a high-strength hot-dip galvanized steel sheet having a hot-dip galvanized layer on the surface of the steel sheet, and a high-strength alloyed hot-dip galvanized sheet having an alloyed hot-dip galvanized layer on the surface of the steel sheet. Steel plates are also included.

The high-strength steel sheet of the present invention has a step of heating to a temperature range of Ac ₁ point + 20 ° C. or higher and Ac ₃ point + 20 ° C. or lower, a step of holding in the temperature range for 50 seconds or longer, and an arbitrary satisfying the following formula (1) The step of cooling at an average cooling rate of 2 to 50 ° C./second up to the temperature T of the step, the step of holding for 10 to 100 seconds in a temperature range satisfying the following formula (1), and the temperature range satisfying the following formula (2) for 200 seconds It can manufacture by the method of including the process hold | maintained in this order.
400 ° C. ≦ T1 (° C.) ≦ 540 ° C. (1)
200 ° C. ≦ T2 (° C.) <400 ° C. (2)

  According to the present invention, after generating polygonal ferrite so that the area ratio with respect to the entire metal structure exceeds 50%, particularly as bainite, there are two types of bainite in which residual γ and carbide are different in form, Bainite generated in a high temperature range of 400 ° C. or higher and 540 ° C. or lower (hereinafter sometimes referred to as high temperature range bainite) and bainite generated in a low temperature range of 200 ° C. or higher and lower than 400 ° C. (hereinafter referred to as low temperature range). In some cases, a high strength steel sheet excellent in workability with good elongation and local deformability can be realized even in a high strength region of 590 MPa or more. Moreover, according to this invention, the manufacturing method of the steel plate which made compatible such high intensity | strength and favorable workability can be provided.

FIG. 1 is a schematic view showing an example of an average interval between adjacent retained austenite and / or carbide. FIG. 2 is a diagram schematically illustrating a distribution state of high-temperature region-generated bainite, low-temperature region-generated bainite, and the like (low-temperature region-generated bainite + tempered martensite). FIG. 3 is a schematic diagram illustrating an example of a heat pattern in the T1 temperature range and the T2 temperature range. FIG. 4 is a graph showing the relationship between tensile strength (TS) and elongation (EL).

The present inventors have repeatedly studied to improve the workability of a high-strength steel sheet having a tensile strength of 590 MPa or more, particularly the elongation and local deformability. as a result,
(1) The metal structure of the steel sheet is mainly composed of polygonal ferrite (specifically, the area ratio with respect to the entire metal structure is more than 50%), and then a mixed structure containing bainite, tempered martensite, and residual γ, Especially as bainite,
(1a) The average interval of the distances between the center positions of adjacent residual γ, adjacent carbides, or adjacent residual γ and adjacent carbides (hereinafter, these may be collectively referred to as residual γ). High temperature region bainite which is 1 μm or more;
(1b) Excellent processability with improved local deformability without deteriorating elongation if two types of bainite, low temperature region bainite, with an average distance between center positions such as residual γ of less than 1 μm are generated. Providing high strength steel sheets
(2) Specifically, the high-temperature region-generated bainite contributes to improvement in elongation of the steel sheet, and the low-temperature region-generated bainite contributes to improvement in local deformability of the steel sheet,
(3) In order to produce a predetermined amount of two types of bainite, after heating in the above two-phase temperature range, any temperature range from 400 ° C. to 540 ° C. (hereinafter sometimes referred to as T1 temperature range). Is cooled at an average cooling rate of 2 ° C./second or more and maintained at this T1 temperature range for 10 to 100 seconds to produce a high temperature range bainite. Hereinafter, it may be referred to as a T2 temperature range), and may be held for 200 seconds or more in this T2 temperature range.
The present invention has been completed.

  First, the metal structure that characterizes the high-strength steel sheet according to the present invention will be described.

《Metallic structure》
The metal structure of the high-strength steel sheet according to the present invention is a mixed structure composed of polygonal ferrite, bainite, tempered martensite, and residual γ.

[Polygonal ferrite]
The metal structure of the steel sheet of the present invention is mainly composed of polygonal ferrite. The main body means that the area ratio with respect to the whole metal structure is more than 50%. Polygonal ferrite is softer than bainite and is a structure that acts to improve the workability by increasing the elongation of the steel sheet. In order to exert such an effect, the area ratio of polygonal ferrite is more than 50%, preferably 55% or more, more preferably 60% or more with respect to the entire metal structure. The upper limit of the area ratio of polygonal ferrite is determined in consideration of the space factor of residual γ measured by the saturation magnetization method, and is, for example, 85%.

  The average equivalent circle diameter D of the polygonal ferrite grains is preferably 10 μm or less (not including 0 μm). By reducing the average equivalent circle diameter D of the polygonal ferrite grains and finely dispersing them, the elongation of the steel sheet can be further improved. Although the detailed mechanism is not clear, by making the polygonal ferrite finer, the dispersion state of the polygonal ferrite with respect to the entire metal structure becomes uniform, so that non-uniform deformation is less likely to occur, which further increases the elongation. It is thought that it contributes to improvement. That is, since the metal structure of the steel sheet of the present invention is composed of a mixed structure of polygonal ferrite, bainite, tempered martensite, and residual γ, when the grain size of polygonal ferrite grains increases, the size of each structure increases. It is considered that since unevenness occurs, non-uniform deformation occurs, strain is concentrated locally, and it becomes difficult to improve workability (particularly, the elongation improving effect due to the formation of polygonal ferrite). Therefore, the average equivalent circle diameter D of polygonal ferrite is preferably 10 μm or less, more preferably 8 μm or less, still more preferably 5 μm or less, and particularly preferably 4 μm or less.

  The area ratio and the average equivalent circle diameter D of the polygonal ferrite can be measured by observing with a scanning electron microscope (SEM).

[Bainite and tempered martensite]
The steel sheet of the present invention is characterized in that the bainite is composed of a composite structure of a high-temperature region-generated bainite and a low-temperature region-generated bainite having a strength higher than that of the high-temperature region-generated bainite. High temperature zone bainite contributes to the improvement of elongation of the steel sheet, and low temperature zone bainite contributes to improvement of local deformability of the steel plate. By including these two types of bainite structures, the local deformability can be improved without deteriorating the elongation of the steel sheet, and the overall workability of the steel sheet can be improved. This is thought to be due to the fact that work hardening ability is increased because non-uniform deformation occurs by combining bainite structures having different strength levels.

The high temperature range bainite is a T1 temperature range of 400 ° C. or more and 540 ° C. or less in the cooling process after heating to a temperature of Ac ₁ point + 20 ° C. or higher and Ac ₃ point + 20 ° C. or lower (two-phase temperature range). It is a bainite structure to be generated. High temperature region bainite is a structure in which the average interval of residual γ and the like is 1 μm or more when a cross section of a steel plate that has undergone nital corrosion is observed by SEM.

  On the other hand, the said low temperature range production | generation bainite is a bainite structure | tissue produced | generated in the T2 temperature range of 200 degreeC or more and less than 400 degreeC in the cooling process after heating to the said two-phase temperature range. Low-temperature region-generated bainite is a structure in which the average interval of residual γ and the like is less than 1 μm when a steel cross section subjected to nital corrosion is observed by SEM.

  Here, the “average interval of residual γ” is the distance between the center positions of adjacent residual γ, the distance between the center positions of adjacent carbides, or adjacent residual γ when the steel sheet cross section is observed by SEM. It is the value which averaged the result of having measured the distance between center positions with a carbide | carbonized_material. The distance between the center positions means a distance between the center positions obtained for each remaining γ or each carbide when measured with respect to the most adjacent residual γ and / or carbide. The center position determines the major axis and minor axis of the residual γ or carbide, and is the position where the major axis and minor axis intersect.

  However, when residual γ or carbide precipitates on the lath boundary, a plurality of residual γ and carbide are connected to form a needle or plate, so the distance between the center positions is the residual γ and / or carbide. Instead of the distance between each other, as shown in FIG. 1, the distance between the lines (the distance between the laths) formed by the residual γ and / or carbides continuously in the major axis direction may be set as the distance between the center positions.

  Moreover, tempered martensite is a structure | tissue which has an effect | action similar to the said low temperature range production | generation bainite, and contributes to the local deformability improvement of a steel plate. Note that the low-temperature region-generated bainite and the tempered martensite cannot be distinguished even by SEM observation. Therefore, in the present invention, the low-temperature region-generated bainite and the tempered martensite are collectively referred to as “low-temperature region-generated bainite and the like”.

  In the present invention, a high-strength steel sheet having improved workability in general can be realized by making the bainite a composite bainite structure including a high-temperature region-generated bainite and a low-temperature region-generated bainite. That is, since the high temperature region generation bainite is softer than the low temperature region generation bainite and the like, it contributes to improving the workability by increasing the elongation (EL) of the steel sheet. On the other hand, low temperature region bainite has low carbides and residual γ, and stress concentration is reduced during deformation. Therefore, the stretch flangeability (λ) and bendability (R) of the steel sheet are improved to improve local deformability. Contributes to improving processability. And in this invention, since such high temperature range production | generation bainite, low temperature range production | generation bainite, etc. are mixed, work hardening ability improves and local deformability can be improved without deteriorating elongation.

  In the present invention, the reason for distinguishing bainite into “high temperature region bainite” and “low temperature region bainite” by the difference in the generation temperature region and the difference in the average interval such as residual γ as described above is a general academic reason. This is because it is difficult to clearly distinguish bainite in the tissue classification. For example, lath-shaped bainite and bainitic ferrite are classified into upper bainite and lower bainite according to the transformation temperature. However, in the steel type containing a large amount of Si of 1.0% or more as in the present invention, precipitation of carbides accompanying the bainite transformation is suppressed, so it is difficult to distinguish these including the martensite structure by SEM observation. It is. Therefore, in the present invention, bainite is not classified based on an academic organization definition, but is distinguished based on the difference in generation temperature range and the average interval such as residual γ as described above.

  The distribution state of the high temperature zone bainite and the low temperature zone bainite is not particularly limited, and both the high temperature zone bainite and the low temperature zone bainite may be generated in the old γ grain, or for each old γ grain A high temperature region generation bainite, a low temperature region generation bainite, or the like may be generated.

  FIG. 2 schematically shows the distribution state of the high temperature region bainite and the low temperature region bainite. In FIG. 2, the high temperature region generation bainite is hatched, and the low temperature region generation bainite is marked with fine dots. FIG. 2 (a) shows a state in which both high-temperature region-generated bainite and low-temperature region-generated bainite are mixed and formed in the old γ grain, and FIG. The high temperature region bainite and the low temperature region bainite are generated. The black circles shown in FIG. 2 indicate the MA mixed phase. The MA mixed phase will be described later.

  In the present invention, when the area ratio of the high temperature region bainite occupying the entire metal structure is b, and the total area ratio of the low temperature region bainite and the like (low temperature region bainite and tempered martensite) occupying the entire metal structure is c, The area ratios b and c must satisfy 5 to 40%. Here, the reason why the total area ratio of the low temperature region-generated bainite and the tempered martensite is defined instead of the area ratio of the low temperature region-generated bainite is that, as described above, these structures cannot be distinguished by SEM observation.

  The area ratio b is 5 to 40%. When there is too little production amount of high temperature range production | generation bainite, the elongation of a steel plate will fall and workability cannot be improved. Therefore, the area ratio b is 5% or more, preferably 8% or more, more preferably 10% or more. However, if the amount of high-temperature region-generated bainite is excessive, the balance of the amount of low-temperature-region-generated bainite and the like deteriorates, and the effect of combining the high-temperature region-generated bainite and the low-temperature region-generated bainite cannot be exhibited. Therefore, the area ratio b of the high temperature region bainite is 40% or less, preferably 35% or less, more preferably 30% or less, and still more preferably 25% or less.

  The total area ratio c is 5 to 40%. If there is too little production amount of low temperature region bainite etc., the local deformability of a steel plate will fall and workability cannot be improved. Therefore, the total area ratio c is 5% or more, preferably 8% or more, more preferably 10% or more. However, if the production amount of low temperature region bainite or the like becomes excessive, the balance of the production amount with the high temperature region bainite is deteriorated, and the effect of combining the low temperature region bainite and the high temperature region bainite is not exhibited. Accordingly, the area ratio c of the low temperature region bainite or the like is 40% or less, preferably 35% or less, more preferably 30% or less, and still more preferably 25% or less.

  The relationship between the area ratio b and the total area ratio c is not particularly limited as long as each range satisfies the above range, and any aspect of b> c, b <c, and b = c is included.

  What is necessary is just to determine the mixing ratio of a high temperature range production | generation bainite, a low temperature range production | generation bainite, etc. according to the characteristic requested | required of a steel plate. Specifically, in order to further improve the local deformability (especially stretch flangeability (λ)) of the workability of the steel sheet, the ratio of the high-temperature region-generated bainite is made as small as possible, and the ratio of the low-temperature region-generated bainite, etc. It should be as large as possible. On the other hand, in order to further improve the elongation of the workability of the steel sheet, the ratio of the high-temperature region-generated bainite should be as large as possible, and the ratio of the low-temperature region-generated bainite should be as small as possible. Further, in order to further increase the strength of the steel sheet, the ratio of the low temperature region bainite or the like may be increased as much as possible, and the ratio of the high temperature region bainite may be decreased as much as possible.

  In the present invention, bainite includes bainitic ferrite. Bainite is a structure in which carbide is precipitated, and bainitic ferrite is a structure in which carbide is not precipitated.

[Polygonal ferrite + bainite + tempered martensite]
In the present invention, the sum of the area ratio a of the polygonal ferrite, the area ratio b of the high-temperature region-generated bainite, and the total area ratio c of the low-temperature region-generated bainite and the like (low-temperature region-generated bainite + tempered martensite) (a + b + c) However, it is preferable that 70% or more of the entire metal structure is satisfied. If the total area ratio (a + b + c) is less than 70%, the elongation may deteriorate. The total area ratio (a + b + c) is more preferably 75% or more, and still more preferably 80% or more. The upper limit of the total area ratio (a + b + c) is determined in consideration of the space factor of residual γ measured by the saturation magnetization method, and is 95%, for example.

[Residual γ]
Residual γ has the effect of accelerating the hardening of the deformed part by transforming into martensite when the steel sheet is deformed under stress, thereby preventing the concentration of strain, thereby improving the uniform deformability and achieving good elongation. Demonstrate. Such an effect is generally called a TRIP effect.

  In order to exert these effects, the volume ratio of the residual γ with respect to the entire metal structure needs to be 5% or more when measured by the saturation magnetization method. The residual γ is preferably 8% by volume or more, more preferably 10% by volume or more. However, if the amount of residual γ generated is too large, the MA mixed phase described later is excessively generated and the MA mixed phase is likely to be coarsened, so that the local deformability (stretch flangeability and bendability) is lowered. Therefore, the upper limit of the residual γ is about 30% by volume, preferably 25% by volume.

  Residual γ is mainly generated between the laths of the metal structure, but as a part of the MA mixed phase, which will be described later, on the aggregate of the lath-like structure (for example, blocks and packets) and the grain boundaries of the old γ May be present in bulk.

[Others]
As described above, the metallographic structure of the steel sheet according to the present invention includes polygonal ferrite, bainite, tempered martensite, and residual γ, and may be composed only of these. As long as it is not impaired, (a) an MA mixed phase in which quenched martensite and residual γ are combined, and (b) a remaining structure such as pearlite may be present.

(A) MA mixed phase The MA mixed phase is generally known as a composite phase of quenched martensite and residual γ, and a part of the structure existing as untransformed austenite before the final cooling is It is a structure formed by transformation into martensite at the time of final cooling, and the rest as austenite. The MA mixed phase thus formed is a very hard structure because carbon is concentrated at a high concentration in the process of heat treatment (especially austempering), and a part thereof has a martensite structure. For this reason, the hardness difference between the bainite and the MA mixed phase is large, and stress is concentrated during deformation, which tends to be a starting point for voids. Therefore, when the MA mixed phase is excessively generated, stretch flangeability and bendability are deteriorated and local deformability is reduced. Decreases. Moreover, when MA mixed phase produces | generates excessively, there exists a tendency for intensity | strength to become high too much. The MA mixed phase is easily generated as the residual γ amount is increased and the Si content is increased. However, the generated amount is preferably as small as possible.

  The MA mixed phase is preferably 30 area% or less, more preferably 25 area% or less, still more preferably 20 area% or less with respect to the entire metal structure when the metal structure is observed with an optical microscope. .

  In the MA mixed phase, the number ratio of MA mixed phases having an equivalent circle diameter d exceeding 7 μm is preferably less than 15% (including 0%) with respect to the total number of MA mixed phases. A coarse MA mixed phase having an equivalent circle diameter d exceeding 7 μm adversely affects local deformability. The ratio of the number of MA mixed phases having the equivalent circle diameter d exceeding 7 μm is preferably less than 10%, more preferably less than 5%, based on the total number of MA mixed phases.

  The number ratio of MA mixed phases having an equivalent circle diameter d exceeding 7 μm may be calculated by observing a cross-sectional surface parallel to the rolling direction with an optical microscope.

  The MA mixed phase is recommended to be as small as possible because experiments have shown that the MA mixed phase tends to generate voids as its particle size increases.

(B) Perlite The pearlite is preferably 20 area% or less with respect to the entire metal structure when the metal structure is observed by SEM. When the area ratio of pearlite exceeds 20%, elongation deteriorates and it becomes difficult to improve workability. The area ratio of pearlite is more preferably 15% or less, further preferably 10% or less, and particularly preferably 5% or less with respect to the entire metal structure.

  The metal structure can be measured by the following procedure.

  Polygonal ferrite, high temperature zone bainite, low temperature zone bainite, etc. (low temperature zone bainite + tempered martensite), and pearlite are subjected to nital corrosion at 1/4 of the thickness of the cross section parallel to the rolling direction of the steel sheet. It can be identified by SEM observation at a magnification of about 3000 times.

  Polygonal ferrite is observed as crystal grains containing no residual γ and carbides observed in white or light gray inside the crystal grains.

  High temperature region bainite and low temperature region bainite are mainly observed in gray, and are observed as a structure in which residual γ and the like observed in white or light gray are dispersed in crystal grains. Therefore, according to SEM observation, since the high temperature region-generated bainite and the low temperature region-generated bainite include residual γ and carbides, the area ratio including the residual γ is calculated.

  Pearlite is observed as a structure in which carbide and ferrite are layered.

  When the cross section of the steel sheet is subjected to Nital corrosion, both carbide and residual γ are observed as a white or light gray structure, and it is difficult to distinguish the two. Among these, carbides (for example, cementite) tend to precipitate in the lath more than between the laths as they are produced in the low temperature range, so when the spacing between the carbides is wide, it is considered that they were produced in the high temperature range, When the interval between the carbides is narrow, it can be considered that the carbides are generated in a low temperature range. Residual γ is usually generated between the laths, but the size of the lath becomes smaller as the tissue generation temperature decreases. Therefore, when the distance between the residual γ is wide, it is considered that the residual γ was generated in a high temperature range. When the interval of is narrow, it can be considered that it was generated in a low temperature region. Therefore, in the present invention, the Nital-corroded cross section is observed with an SEM, focusing on residual γ observed as white or light gray in the observation field and measuring the distance between the center positions of adjacent residual γ, etc. A structure having an average value (average interval) of 1 μm or more is referred to as a high-temperature region generation bainite, and a structure having an average interval of less than 1 μm is referred to as a low-temperature region generation bainite.

  Since the residual γ cannot be identified by SEM observation, the volume fraction is measured by the saturation magnetization method. This volume ratio value can be read as the area ratio as it is. The detailed measurement principle by the saturation magnetization method may be referred to “R & D Kobe Steel Engineering Reports, Vol.52, No.3, 2002, p.43-46”.

  Thus, while the volume ratio (area ratio) of residual γ is measured by the saturation magnetization method, the area ratio of high temperature region bainite and low temperature region bainite is measured by SEM observation including residual γ. Therefore, the sum of these may exceed 100%.

  The MA mixed phase is observed as a white structure when subjected to repeller corrosion at a 1/4 position of the plate thickness in a cross section parallel to the rolling direction of the steel plate and observed with an optical microscope at a magnification of about 1000 times.

  Next, the chemical component composition of the high-strength steel sheet according to the present invention will be described.

<About component composition>
The high-strength steel sheet of the present invention contains C: 0.10 to 0.3%, Si: 1.0 to 3%, Mn: 1.0 to 2.5%, Al: 0.005 to 3%. And P: 0.1% or less (not including 0%) and S: 0.05% or less (not including 0%). The reason for setting this range is as follows.

  C is an element necessary for increasing the strength of the steel sheet and generating residual γ. Therefore, the amount of C is 0.10% or more, preferably 0.13% or more, more preferably 0.15% or more. However, when C is contained excessively, weldability is lowered. Therefore, the C content is 0.3% or less, preferably 0.25% or less, more preferably 0.20% or less.

  Si contributes to increasing the strength of the steel sheet as a solid solution strengthening element, and suppresses the precipitation of carbide during holding in the T1 temperature range and T2 temperature range described later (during the austempering process), thereby reducing the residual γ. It is an extremely important element for effective generation. Accordingly, the Si amount is 1.0% or more, preferably 1.2% or more, more preferably 1.3% or more. However, when Si is excessively contained, reverse transformation to the γ phase does not occur during heating and soaking in annealing, and a large amount of polygonal ferrite remains, resulting in insufficient strength. In addition, Si scale is generated on the surface of the steel sheet during hot rolling to deteriorate the surface properties of the steel sheet. Therefore, the amount of Si is 3% or less, preferably 2.50% or less, more preferably 2.0% or less.

  Mn is an element necessary for obtaining bainite and tempered martensite. Mn is an element that effectively acts to stabilize γ and generate residual γ. In order to exert such an effect, the amount of Mn is set to 1.0% or more, preferably 1.5% or more, more preferably 1.8% or more. However, when Mn is contained excessively, the generation of high temperature region bainite is remarkably suppressed. Further, excessive addition of Mn causes deterioration of weldability and workability due to segregation. Therefore, the Mn content is 2.5% or less, preferably 2.4% or less, and more preferably 2.3% or less.

  Al, like Si, is an element that suppresses the precipitation of carbides during the austempering process and contributes to the formation of residual γ. Al is an element that acts as a deoxidizer in the steel making process. Therefore, the Al content is 0.005% or more, preferably 0.01% or more, more preferably 0.03% or more. However, when Al is contained excessively, the inclusions in the steel sheet increase so much that ductility deteriorates. Therefore, the Al content is 3% or less, preferably 1.5% or less, more preferably 1% or less, and still more preferably 0.5% or less.

  P is an impurity element inevitably contained in the steel. When the amount of P is excessive, the weldability of the steel sheet is deteriorated. Therefore, the amount of P is 0.1% or less, preferably 0.08% or less, more preferably 0.05% or less. The amount of P is preferably as small as possible, but it is industrially difficult to reduce it to 0%.

  S is an impurity element inevitably contained in the steel, and is an element that deteriorates the weldability of the steel sheet, as in the case of P. Further, S forms sulfide-based inclusions in the steel sheet, and when this increases, the workability decreases. Therefore, the amount of S is 0.05% or less, preferably 0.01% or less, more preferably 0.005% or less. The amount of S should be as small as possible, but it is industrially difficult to make it 0%.

  The high-strength steel sheet according to the present invention satisfies the above component composition, and the remaining components are iron and inevitable impurities other than P and S. Examples of inevitable impurities include N, O (oxygen), and trump elements (eg, Pb, Bi, Sb, Sn, etc.). Of the inevitable impurities, the N content is preferably 0.01% or less (not including 0%), and the O content is preferably 0.01% or less (not including 0%).

  N is an element that contributes to strengthening of the steel sheet by precipitating nitrides in the steel sheet. However, when N is excessively contained, a large amount of nitride precipitates and causes elongation, stretch flangeability, and deterioration of bendability. cause. Accordingly, the N content is preferably 0.01% or less, more preferably 0.008% or less, and still more preferably 0.005% or less.

  O (oxygen) is an element that causes a decrease in elongation, stretch flangeability, and bendability when contained in excess. Therefore, the O content is preferably 0.01% or less, more preferably 0.005% or less, and still more preferably 0.003% or less.

The steel sheet of the present invention is further as another element,
(A) Cr: 1% or less (not including 0%) and / or Mo: 1% or less (not including 0%),
(B) A group consisting of Ti: 0.15% or less (not including 0%), Nb: 0.15% or less (not including 0%), and V: 0.15% or less (not including 0%) One or more elements selected from
(C) Cu: 1% or less (not including 0%) and / or Ni: 1% or less (not including 0%),
(D) B: 0.005% or less (excluding 0%),
(E) Ca: 0.01% or less (not including 0%), Mg: 0.01% or less (not including 0%), and rare earth elements: 0.01% or less (not including 0%) One or more elements selected from the group,
Etc. may be contained.

  (A) Similar to Mn, Cr and Mo are elements that act effectively to obtain bainite and tempered martensite. These elements can be used alone or in combination. In order to effectively exhibit such an action, Cr and Mo are each preferably contained alone in an amount of 0.1% or more, more preferably 0.2% or more. However, when the content of Cr and Mo exceeds 1%, the generation of high temperature region bainite is remarkably suppressed. In addition, excessive addition increases the cost. Accordingly, Cr and Mo are each preferably 1% or less, more preferably 0.8% or less, and still more preferably 0.5% or less. When Cr and Mo are used in combination, the total amount is recommended to be 1.5% or less.

  (B) Ti, Nb, and V are elements that form precipitates such as carbides and nitrides in the steel sheet, strengthen the steel sheet, and also have the effect of refining the polygonal ferrite grains by refining the old γ grains. is there. In order to exhibit such an action effectively, Ti, Nb and V are each preferably contained in an amount of 0.01% or more, more preferably 0.02% or more. However, when it contains excessively, carbide will precipitate to a grain boundary and the stretch flangeability and bendability of a steel plate will deteriorate. Therefore, Ti, Nb and V are each independently preferably 0.15% or less, more preferably 0.12% or less, and still more preferably 0.1% or less. Ti, Nb, and V may each be contained alone, or may contain two or more elements that are arbitrarily selected.

  (C) Cu and Ni are elements that effectively act to stabilize γ and generate residual γ. These elements can be used alone or in combination. In order to exhibit such an action effectively, it is preferable to contain Cu and Ni individually by 0.05% or more, more preferably 0.1% or more. However, when Cu and Ni are contained excessively, the hot workability deteriorates. Accordingly, Cu and Ni are each preferably preferably 1% or less, more preferably 0.8% or less, and still more preferably 0.5% or less. In addition, when Cu is contained in excess of 1%, hot workability deteriorates. However, when Ni is added, deterioration of hot workability is suppressed. However, Cu may be added in excess of 1%.

  (D) B is an element that acts effectively to produce bainite and tempered martensite, as in the case of Mn, Cr and Mo. In order to exhibit such an action effectively, B is preferably contained in an amount of 0.0005% or more, more preferably 0.001% or more. However, when B is contained excessively, a boride is generated in the steel sheet and the ductility is deteriorated. Moreover, when B is contained excessively, the production | generation of high temperature range production | generation bainite will be suppressed remarkably similarly to said Cr and Mo. Accordingly, the B content is preferably 0.005% or less, more preferably 0.004% or less, and still more preferably 0.003% or less.

  (E) Ca, Mg and rare earth elements (REM) are elements that act to finely disperse inclusions in the steel sheet. In order to exhibit such an action effectively, Ca, Mg and rare earth elements are each preferably contained in an amount of 0.0005% or more, more preferably 0.001% or more. However, when it contains excessively, castability, hot workability, etc. will deteriorate and it will become difficult to manufacture. Further, excessive addition causes the ductility of the steel sheet to deteriorate. Therefore, Ca, Mg, and rare earth elements are each independently preferably 0.01% or less, more preferably 0.005% or less, and still more preferably 0.003% or less.

  The rare earth element means a lanthanoid element (15 elements from La to Lu), Sc (scandium) and Y (yttrium), and among these elements, it is selected from the group consisting of La, Ce and Y. It is preferable to contain at least one kind of element, more preferably La and / or Ce.

  The high-strength steel sheet according to the present invention is excellent in workability because it has a tensile strength of 590 MPa or more, excellent elongation, and good local deformability. This high-strength steel plate is suitably used as a material for structural parts of automobiles. Structural parts of automobiles include, for example, front and rear side members and crashing parts such as crash boxes, reinforcing materials such as pillars (for example, center pillar reinforcement), roof rail reinforcing materials, side sills, floor members, Examples include vehicle body components such as kick parts, shock-absorbing parts such as bumper reinforcements and door impact beams, and seat parts.

  The high-strength steel sheet can be suitably used as a material for warm forming because of its good workability in warm conditions. In addition, warm processing means shape | molding in the temperature range of about 50-500 degreeC.

  The metal structure and component composition of the high-strength steel sheet according to the present invention have been described above.

Next, a method capable of producing the high-strength steel plate will be described. The high-strength steel sheet has a step of heating a steel sheet satisfying the above component composition to a temperature range (two-phase temperature range) of Ac ₁ point + 20 ° C. or higher and Ac ₃ point + 20 ° C. or lower, and hold at the temperature range for 50 seconds or longer. A step of cooling to an arbitrary temperature T satisfying the following formula (1) at an average cooling rate of 2 to 50 ° C./second, a step of holding for 10 to 100 seconds in a temperature range satisfying the following formula (1), It can manufacture by including in the order the process of hold | maintaining for 200 second or more in the temperature range which satisfy | fills following formula (2). Hereinafter, each step will be described in order.
400 ° C. ≦ T1 (° C.) ≦ 540 ° C. (1)
200 ° C. ≦ T2 (° C.) <400 ° C. (2)

First, a hot-rolled steel sheet obtained by hot rolling a slab according to a conventional method as a high-strength steel sheet before heating to a two-phase temperature range [temperature range of Ac ₁ point + 20 ° C. or higher, Ac ₃ point + 20 ° C. or lower] Prepare a cold rolled product. In hot rolling, the finish rolling temperature may be set to, for example, 800 ° C. or more, and the winding temperature may be set to, for example, 700 ° C. or less. In cold rolling, the cold rolling rate may be rolled in a range of 10 to 70%, for example.

The cold-rolled steel sheet obtained by cold rolling is heated to a temperature range of Ac ₁ point + 20 ° C. or higher and Ac ₃ point + 20 ° C. or lower in a continuous annealing line, and kept at this temperature range for 50 seconds or more and soaking. To do.

By setting the heating temperature to a two-phase temperature range of ferrite and austenite, a predetermined amount of polygonal ferrite can be generated. That is, if the heating temperature is too high, an austenite single-phase region is formed, and the formation of polygonal ferrite is suppressed, so that the elongation of the steel sheet cannot be improved and the workability deteriorates. Accordingly, the heating temperature is Ac ₃ point + 20 ° C. or less, preferably Ac ₃ point + 10 ° C. or less, more preferably less than Ac ₃ point. In addition, when heated to Ac ₃ point or higher, it becomes the temperature range of the austenite single phase, but if the heating temperature is Ac ₃ point + 20 ° C. or less in the soaking time specified in the present invention, the soaking is maintained. Since a small amount of polygonal ferrite remains even if it is carried out, a predetermined amount of polygonal ferrite can be generated by adjusting the average cooling rate after soaking as described later. However, if the heating temperature is less than Ac ₁ point + 20 ° C., the amount of polygonal ferrite produced becomes excessive, and a predetermined amount of high temperature region bainite, low temperature region bainite, etc. and residual γ cannot be obtained, so workability is improved. to degrade. Accordingly, the heating temperature is Ac ₁ point + 20 ° C. or higher, preferably Ac ₁ point + 30 ° C. or higher, more preferably Ac ₁ point + 50 ° C. or higher.

  If the soaking time in the two-phase temperature range is less than 50 seconds, the steel sheet cannot be heated uniformly, so that the formation of residual γ is suppressed, elongation and local deformability are reduced, and workability cannot be improved. Therefore, the soaking time is 50 seconds or longer, preferably 100 seconds or longer. However, if the soaking time is too long, the austenite grain size becomes large, and the polygonal ferrite grains are coarsened accordingly, and the elongation and local deformability tend to deteriorate. Therefore, the soaking time is preferably 500 seconds or shorter, more preferably 450 seconds or shorter.

  In addition, what is necessary is just to let the average heating rate when heating the said cold-rolled steel plate to the said two-phase temperature range be 1 degree-C / sec or more, for example.

The Ac ₁ point and Ac ₃ point are calculated from the following formulas (a) and (b) described in “Leslie Steel Materials Science” (Maruzen Co., Ltd., issued May 31, 1985, P.273). it can. In the following formulas (a) and (b), [] indicates the content (mass%) of each element, and the content of elements not included in the steel sheet may be calculated as 0 mass%.
Ac ₁ (° C.) = 723-10.7 × [Mn] −16.9 × [Ni] + 29.1 × [Si] + 16.9 × [Cr] (a)
Ac ₃ (° C.) = 910−203 × [C] ^1/2 + 44.7 × [Si] −30 × [Mn] −11 × [Cr] + 31.5 × [Mo] −20 × [Cu] −15 2 × [Ni] + 400 × [Ti] + 104 × [V] + 700 × [P] + 400 × [Al] (b)

  After heating to the above two-phase temperature range and holding for 50 seconds or more and soaking, it is cooled at an average cooling rate of 2 to 50 ° C./second to any temperature T satisfying the above formula (1). By cooling the range from the two-phase temperature range to an arbitrary temperature T satisfying the above formula (1) at a predetermined average cooling rate or higher, a predetermined amount of polygonal ferrite can be generated. And low temperature region bainite can be generated. When the average cooling rate in this temperature range is less than 2 ° C./second, pearlite transformation occurs, pearlite is excessively generated, elongation decreases, and workability deteriorates. The average cooling rate in this section is preferably 5 ° C./second or more, more preferably 10 ° C./second or more. However, if the average cooling rate in the section is too large, a predetermined amount of polygonal ferrite cannot be secured. Therefore, the average cooling rate is 50 ° C./second or less, preferably 40 ° C./second or less, more preferably 30 ° C./second or less.

  After cooling to an arbitrary temperature T satisfying the above formula (1), after holding for 10 to 100 seconds in the T1 temperature range satisfying the above formula (1), it is performed for 200 seconds or more in the T2 temperature range satisfying the above formula (2). Hold. By appropriately controlling the time for holding in the T1 temperature range and the T2 temperature range, it is possible to generate a predetermined amount of high temperature region bainite, low temperature region bainite, and the like. Specifically, the amount of high-temperature region-generated bainite can be controlled by holding in the T1 temperature region for a predetermined time, and the untransformed austenite is converted into low-temperature region-generated bainite or martensite by the austempering process that is maintained in the T2 temperature region for a predetermined time. While transforming into sites, carbon can be concentrated to austenite to generate residual γ, and a metal structure defined in the present invention can be generated.

  Moreover, the effect which can suppress the production | generation of MA mixed phase is exhibited by combining the holding | maintenance in T1 temperature range, and the holding | maintenance in T2 temperature range. This mechanism is considered as follows. In general, when Si or Al is added, precipitation of carbides is suppressed, so that free carbon exists in the steel, and in the austempering process, carbon is concentrated to untransformed austenite along with bainite transformation. Is recognized. A large amount of residual γ can be generated by concentrating carbon to untransformed austenite.

  Here, a phenomenon in which carbon is concentrated to untransformed austenite will be described. It is known that the bainite transformation also stops because the concentration of carbon is limited to the concentration indicated by the To line where the free energy of ferrite and austenite becomes equal. Since the To line becomes lower in carbon concentration as the temperature is higher, if the austempering process is performed at a relatively high temperature, the bainite transformation stops at a certain level even if the processing time is increased. At this time, since the stability of untransformed austenite is low, a coarse MA mixed phase is generated.

  Therefore, in the present invention, after holding in the T1 temperature range, the allowable amount of C concentration to the untransformed austenite can be increased by holding in the T2 temperature range, so the low temperature range is higher than the high temperature range. The bainite transformation proceeds and the MA mixed phase becomes smaller. Further, since the size of the lath-like structure is smaller when held at the T2 temperature range than when held at the T1 temperature range, the MA mixed phase itself is subdivided even if the MA mixed phase exists. Thus, the MA mixed phase can be reduced. Furthermore, since it hold | maintains in T2 temperature range after hold | maintaining for a predetermined time in T1 temperature range, the high temperature range production | generation bainite has already produced | generated when the holding | maintenance in T2 temperature range was started. Accordingly, in the T2 temperature region, the high temperature region-generated bainite is a trigger, and the transformation of the low temperature region-generated bainite is promoted, so that the effect of shortening the time of the austempering treatment is also exhibited.

  In the case where the temperature is cooled from the two-phase temperature range to an arbitrary temperature satisfying the formula (2) without being held in the T1 temperature range, and held only in the T2 temperature range satisfying the formula (2). Even in the simple austempering process at a low temperature, the size of the lath-like structure is reduced, so that the MA mixed phase can be reduced. However, in this case, since the temperature is not maintained in the T1 temperature range, almost no high temperature range bainite is generated, the dislocation density of the base lath structure is increased, elongation and local deformability are reduced, and workability is reduced. Deteriorates.

  In the present invention, the T1 temperature range defined by the above formula (1) is specifically 400 ° C. or higher and 540 ° C. or lower. By maintaining for a predetermined time in this temperature range, high temperature range bainite can be generated. That is, when the temperature is maintained at a temperature exceeding 540 ° C., the formation of high temperature bainite is suppressed. On the other hand, polygonal ferrite is excessively generated and pseudo pearlite is generated, so that desired characteristics cannot be obtained. Therefore, the upper limit of the T1 temperature range is 540 ° C, preferably 520 ° C, more preferably 500 ° C. On the other hand, if the holding temperature is lower than 400 ° C., high temperature region bainite is not generated, so that elongation is lowered and workability cannot be improved. Therefore, the lower limit of the T1 temperature range is 400 ° C, preferably 420 ° C.

  The time held in the T1 temperature range is 10 to 100 seconds. If the holding time exceeds 100 seconds, the high-temperature region-generated bainite is excessively generated. Therefore, as will be described later, the amount of low-temperature region-generated bainite or the like cannot be ensured even if the predetermined time is maintained in the T2 temperature region. Accordingly, it is impossible to achieve both strength and workability. Further, if the temperature is held for a long time in the T1 temperature range, carbon is excessively concentrated in the austenite, so that a coarse MA mixed phase is generated even if austempering is performed in the T2 temperature range, and workability deteriorates. Therefore, the holding time is 100 seconds or less, preferably 90 seconds or less, more preferably 80 seconds or less. However, if the holding time in the T1 temperature region is too short, the amount of high-temperature region-generated bainite is reduced, so that elongation is lowered and workability cannot be improved. Accordingly, the holding time in the T1 temperature range is 10 seconds or longer, preferably 15 seconds or longer, more preferably 20 seconds or longer, and even more preferably 30 seconds or longer.

  In the present invention, the holding time in the T1 temperature range means the time from when the surface temperature of the steel sheet reaches the upper limit temperature in the T1 temperature range until it reaches the lower limit temperature in the T1 temperature range. That is, it is the time from when the surface temperature of the steel sheet reaches 540 ° C. until it reaches 400 ° C.

  In order to maintain the temperature in the T1 temperature range that satisfies the above formula (1), for example, a heat pattern shown in (i) to (iii) of FIG. 3 may be employed.

  FIG. 3 (i) shows an example in which the temperature is rapidly cooled from the two-phase temperature range to an arbitrary temperature T satisfying the above formula (1), and then kept at this temperature T for a predetermined time. It is cooled to any temperature that satisfies. Although FIG. 3 (i) shows a case where one-stage constant temperature holding is performed, the present invention is not limited to this, and two or more constant temperature holdings having different holding temperatures are performed as long as they are within the T1 temperature range. May be.

  FIG. 3 (ii) shows that after rapidly cooling from the two-phase temperature range to an arbitrary temperature T satisfying the above formula (1), the cooling rate is changed, and after cooling for a predetermined time within the range of the T1 temperature range, In this example, the cooling rate is changed and the cooling is performed to an arbitrary temperature satisfying the above formula (2). FIG. 3 (ii) shows the case of cooling over a predetermined time within the range of the T1 temperature range, but is not limited to this, and the step of heating over the predetermined time if within the T1 temperature range. May be included, and cooling and heating may be repeated as appropriate. Further, as shown in FIG. 3 (ii), not only one-stage cooling but also two-stage or more multi-stage cooling with different cooling rates may be performed. Further, one-stage heating or multi-stage heating of two or more stages may be performed (not shown).

  FIG. 3 (iii) shows that after rapidly cooling from the two-phase temperature range to an arbitrary temperature T satisfying the above formula (1), the cooling rate is changed and the same cooling is performed until an arbitrary temperature satisfying the above formula (2). This is an example of slow cooling at a speed. Even in such a case of slow cooling, the residence time in the T1 temperature range may be 10 to 100 seconds.

  The present invention is not intended to be limited to the heat patterns shown in (i) to (iii) of FIG. 3, and heat patterns other than those described above can be appropriately employed as long as the requirements of the present invention are satisfied.

  In the present invention, the T2 temperature range defined by the above formula (2) is specifically 200 ° C. or more and less than 400 ° C. By maintaining in this temperature range for a predetermined time, untransformed austenite that has not been transformed in the T1 temperature range can be transformed into low temperature range bainite or martensite. Further, by securing a sufficient holding time, the bainite transformation proceeds, finally residual γ is generated, and the MA mixed phase is subdivided. Although this martensite exists as quenching martensite immediately after transformation, it is tempered while being maintained in the T2 temperature region and remains as tempered martensite. This tempered martensite exhibits the same characteristics as low temperature region bainite generated in the temperature region where martensitic transformation occurs. However, if the temperature is maintained at 400 ° C. or higher, a coarse MA mixed phase is generated, so that elongation and local deformability are lowered, and workability cannot be improved. Accordingly, the T2 temperature range is less than 400 ° C., preferably 390 ° C. or less, more preferably 380 ° C. or less. On the other hand, since the low temperature region bainite is not generated even if the temperature is kept below 200 ° C., the carbon concentration in γ is low, the amount of residual γ cannot be secured, and more quenching martensite is generated, so the strength is high. It becomes high and elongation and local deformability deteriorate. Further, since the carbon concentration in γ becomes low and the amount of residual γ cannot be secured, the elongation cannot be increased. Therefore, the lower limit of the T2 temperature range is 200 ° C, preferably 250 ° C, more preferably 280 ° C.

  The time for holding in the T2 temperature range that satisfies the above formula (2) is 200 seconds or longer. If the holding time is less than 200 seconds, the amount of low temperature region bainite and the like is reduced, the carbon concentration in γ is lowered, the amount of residual γ cannot be secured, and more hardened martensite is generated, so the strength is increased. It becomes high and elongation and local deformability deteriorate. Further, since carbon concentration is not promoted, the amount of residual γ is reduced, and the elongation cannot be improved. Moreover, since the MA mixed phase produced | generated in the said T1 temperature range cannot be refined | miniaturized, local deformability cannot be improved. Accordingly, the holding time is 200 seconds or longer, preferably 250 seconds or longer, more preferably 300 seconds or longer. Although the upper limit of the holding time is not particularly limited, productivity decreases when held for a long time, and concentrated carbon cannot be precipitated as carbides to generate residual γ, resulting in a decrease in elongation and workability. to degrade. Therefore, the upper limit of the holding time may be 1800 seconds, for example.

  In the present invention, the holding time in the T2 temperature range means the time from when the surface temperature of the steel sheet reaches the upper limit temperature in the T2 temperature range until it reaches the lower limit temperature in the T2 temperature range. That is, it is the time from reaching the temperature of less than 400 ° C. until reaching 200 ° C.

  The method of holding in the T2 temperature range is not particularly limited as long as the residence time in the T2 temperature range is 200 seconds or more, and may be held at a constant temperature as in the heat pattern in the T1 temperature range, or the T2 temperature. It may be cooled or heated in the zone. Further, multistage holding may be performed at different holding temperatures.

  After holding for a predetermined time in the T2 temperature range, the high strength steel sheet according to the present invention can be manufactured by cooling to room temperature.

  A hot-dip galvanized layer or an alloyed hot-dip galvanized layer may be formed on the surface of the high-strength steel plate.

  Conditions for forming the hot dip galvanized layer or the alloyed hot dip galvanized layer are not particularly limited, and known conditions can be employed.

  For example, the hot dip galvanized layer is preferably formed at a plating bath temperature of 400 to 500 ° C, more preferably 440 to 470 ° C. The composition of the plating bath is not particularly limited, and a known hot dip galvanizing bath may be used.

  An alloyed hot-dip galvanized steel sheet can be produced by subjecting the hot-dip galvanized steel sheet on which the hot-dip galvanized layer is formed to a conventional alloying treatment. The alloying treatment may be performed, for example, at about 450 to 600 ° C. (particularly about 480 to 570 ° C.) and held for about 5 to 30 seconds (particularly about 10 to 25 seconds). The alloying process may be performed using, for example, a heating furnace, a direct fire, or an infrared heating furnace. The heating means is not particularly limited, and for example, conventional means such as gas heating, induction heater heating (heating by a high frequency induction heating device) can be adopted.

  The technique of the present invention can be suitably used particularly for a thin steel plate having a thickness of 3 mm or less.

  Hereinafter, the present invention will be described in more detail with reference to examples. However, the present invention is not limited by the following examples, and is implemented with appropriate modifications within a range that can meet the purpose described above and below. Of course, any of these is also included in the technical scope of the present invention.

  Experimental slabs were manufactured by vacuum melting steel having the chemical composition shown in Table 1 below (the balance being iron and inevitable impurities other than P, S, N, and O). In Table 1 below, REM used Misch metal containing about 50% La and about 30% Ce.

The Ac ₁ point was calculated based on the chemical components shown in Table 1 below and the above formula (a), and the Ac ₃ point was calculated based on the above formula (b). The results are shown in Tables 2 and 3 below.

  The obtained experimental slab was hot-rolled, cold-rolled, and then continuously annealed to produce a specimen. Specific conditions are as follows.

  The experimental slab was heated and held at 1250 ° C. for 30 minutes, then hot rolled so that the reduction rate was about 90% and the final rolling temperature was 920 ° C., and wound at this temperature at an average cooling rate of 30 ° C./second. It cooled to the temperature of 500 degreeC and wound up. After winding, it was kept at this winding temperature (500 ° C.) for 30 minutes, and then cooled to room temperature to produce a hot rolled steel sheet having a thickness of 2.6 mm.

  The obtained hot-rolled steel sheet was pickled to remove the surface scale, and then cold-rolled at a cold rolling rate of 46% to produce a cold-rolled steel sheet having a thickness of 1.4 mm.

  The obtained cold-rolled steel sheet was heated to the temperature (° C.) shown in the following Table 2 and Table 3, held for the time shown in the following Table 2 and Table 3, and soaked, and then any one of the following four patterns The sample was cooled and continuously annealed to produce a specimen.

(Cooling pattern i; corresponding to (i) in FIG. 3)
After soaking, the sample is cooled to the starting temperature T (° C.) shown in Table 2 and Table 3 at the average cooling rate (° C./second) shown in Table 2 and Table 3 below, and then held at this starting temperature T (° C.). Then, it was cooled to the starting temperature (° C.) in the T2 temperature range shown in Tables 2 and 3 below, and kept at this starting temperature. Tables 2 and 3 below show the stay time (seconds) in the T1 temperature range and the stay time (seconds) in the T2 temperature range. In addition, the time (seconds) from the time when the holding is completed in the T1 temperature range until the start temperature in the T2 temperature range is reached is shown.

(Cooling pattern ii; corresponding to (ii) in FIG. 3)
After soaking, after cooling to the start temperature T (° C.) shown in Table 2 and Table 3 at the average cooling rate (° C./second) shown in Table 2 and Table 3 below, the end temperature shown in Table 2 and Table 3 below is shown. It was cooled to (° C.), then cooled to the starting temperature (° C.) in the T2 temperature range shown in Tables 2 and 3 below, and held at this starting temperature for the time (seconds) shown in Tables 2 and 3 below. Tables 2 and 3 below show the stay time (seconds) in the T1 temperature range and the stay time (seconds) in the T2 temperature range. In addition, the time (seconds) from the time when the holding is completed in the T1 temperature range until the start temperature in the T2 temperature range is reached is shown.

(Cooling pattern iii; corresponding to (iii) of FIG. 3 above)
After soaking, it was cooled to the starting temperature T (° C.) shown in Table 2 and Table 3 at the average cooling rate (° C./second) shown in Table 2 and Table 3 below, and then the T2 temperature shown in Table 2 and Table 3 below. Cooled to the starting temperature (° C.) in the zone and held at this starting temperature. Tables 2 and 3 below show the stay time (seconds) in the T1 temperature range and the stay time (seconds) in the T2 temperature range.

(Cooling pattern iv)
After soaking, it was cooled to the starting temperature (° C.) in the T1 temperature range shown in Table 2 below, and kept at this starting temperature. That is, No. 2 in Table 2 below. 19 is an example in which after soaking, the sample was held at 420 ° C. for 450 seconds and then cooled at a time without holding to room temperature (average cooling rate was 5 ° C./second), and stay in the T2 temperature range shown in Table 2 below The time indicates the time required to pass through the T2 temperature range. Table 2 below shows the stay time (seconds) in the T1 temperature range and the stay time (seconds) in the T2 temperature range.

  Of the start temperature and end temperature in the T1 temperature range shown in Table 2, and the start temperature in the T2 temperature range, the values marked with * are outside the T1 temperature range or T2 temperature range defined in the present invention. However, for convenience of explanation, the temperature is described in each column in order to show the heat pattern.

  With respect to the obtained specimens, observation of the metal structure and evaluation of mechanical properties were carried out according to the following procedure.

《Observation of metal structure》
Among metal structures, the area ratio of polygonal ferrite, high temperature region bainite, low temperature region bainite, etc. (ie, low temperature region bainite + tempered martensite) is calculated based on the results of observation with a scanning electron microscope (SEM). The volume fraction of residual γ was measured by the saturation magnetization method.

[(1) Area ratios of polygonal ferrite, high-temperature region-generated bainite, low-temperature region-generated bainite, etc.]
About the cross section parallel to the rolling direction of the test material, the surface was polished, further electropolished, and then subjected to nital corrosion, and the 1/4 position of the plate thickness was observed with SEM at 5 magnifications at 3000 magnifications. The observation visual field was about 50 μm × about 50 μm.

  Next, the average interval between residual γ and carbides observed as white or light gray in the observation field was measured based on the method described above. The area ratios of the high-temperature region-generated bainite and the low-temperature region-generated bainite, which are distinguished by these average intervals, were measured by a point calculation method.

  Tables 4 and 5 below show the area ratio a (%) of polygonal ferrite, the area ratio b (%) of high-temperature region-generated bainite, and the total area ratio c (%) of low-temperature region-generated bainite and tempered martensite. Further, the total area ratio (a + b + c) of the area ratio a, the area ratio b, and the total area ratio c is also shown.

  Moreover, the circle equivalent diameter of the polygonal ferrite grains observed in the observation field was measured, and the average value was obtained. The results are shown in Tables 4 and 5 below. Further, the case where the average equivalent circle diameter D of the polygonal ferrite grains is 10 μm or less is evaluated as ◯, and the case where it exceeds 10 μm is evaluated as Δ, and the evaluation results are shown in Tables 4 and 5 below.

[(2) Volume ratio of residual γ]
Of the metal structure, the volume fraction of residual γ was measured by the saturation magnetization method. Specifically, the saturation magnetization (I) of the specimen and the saturation magnetization (Is) of a standard sample heat-treated at 400 ° C. for 15 hours were measured, and the volume fraction (Vγr) of residual γ was obtained from the following formula. The saturation magnetization was measured at room temperature using a direct-current magnetization BH characteristic automatic recording device “model BHS-40” manufactured by Riken Electronics Co., Ltd. with a maximum applied magnetization of 5000 (Oe).
Vγr = (1−I / Is) × 100

  Moreover, the surface of the cross section parallel to the rolling direction of the test material is polished, and observed with five optical fields at an observation magnification of 1000 times using an optical microscope, corresponding to a circle of MA mixed phase in which residual γ and quenching martensite are combined. The diameter d was measured. The ratio of the number of MA mixed phases in which the equivalent circle diameter d in the observation cross section exceeds 7 μm was calculated with respect to the total number of MA mixed phases. The evaluation results are shown in Tables 4 and 5 below, assuming that the number ratio is less than 15% as pass (◯) and 15% or more as reject (x).

<< Evaluation of mechanical properties >>
The mechanical properties of the specimens were evaluated based on tensile strength (TS), elongation (EL), hole expansion rate (λ), critical bending radius (R), and Erichsen value.

  (1) Tensile strength (TS) and elongation (EL) were measured by conducting a tensile test based on JIS Z2241. The test piece was obtained by cutting a No. 5 test piece defined in JIS Z2201 from the test material so that the direction perpendicular to the rolling direction of the test material was the longitudinal direction. The measurement results are shown in Tables 4 and 5 below.

  (2) Stretch flangeability is evaluated by the hole expansion rate. The hole expansion rate (λ) was measured by performing a hole expansion test based on the Steel Federation Standard JFST 1001. The measurement results are shown in Tables 4 and 5 below.

  (3) The critical bending radius (R) was measured by performing a V-bending test based on JIS Z2248. The test piece is a No. 1 test piece (sheet thickness: 1.4 mm) defined in JIS Z2204 so that the direction perpendicular to the rolling direction of the test material is the longitudinal direction (the bending ridge line coincides with the rolling direction). ) Was cut out from the test material. The V-bending test was performed after mechanical grinding was performed on the end face in the longitudinal direction of the test piece so as not to cause cracks.

  The angle between the die and the punch is 90 °, the tip radius of the punch is changed in units of 0.5 mm, a V-bending test is performed, and the radius of the tip of the punch that can be bent without cracks is determined as the limit bending radius (R). It was. The measurement results are shown in Tables 4 and 5 below. In addition, the presence or absence of crack generation was observed using a loupe, and the determination was made based on the absence of hair crack generation.

  (4) The Eriksen value was measured by conducting an Eriksen test based on JIS Z2247. The test piece used was cut from the test material so as to be 90 mm × 90 mm × 1.4 mm in thickness. The Eriksen test was performed using a punch having a diameter of 20 mm. The measurement results are shown in Tables 4 and 5 below. In addition, according to the Erichsen test, the composite effect by both the total elongation characteristic and local ductility of a steel plate can be evaluated.

  The mechanical properties of the specimens were evaluated according to the criteria of elongation (EL), hole expansion rate (λ), critical bending radius (R), and Erichsen value according to tensile strength (TS). That is, since EL, λ, R, and Erichsen values required by steel sheet TS differ, mechanical characteristics were evaluated according to the following criteria according to the TS level.

  Based on the following evaluation criteria, the case where all the characteristics of EL, λ, R, and Erichsen values are satisfied is accepted (◯), and the case where any characteristic is less than the reference value is rejected (×). The evaluation results are shown in Tables 4 and 5 below.

(1) In the case of 590 MPa class TS: 590 MPa or more and less than 780 MPa EL: 34% or more λ: 30% or more R: 0.5 mm or less Erichsen value: 10.8 mm or more

(2) In the case of 780 MPa class TS: 780 MPa or more and less than 980 MPa EL: 25% or more λ: 30% or more R: 1.0 mm or less Erichsen value: 10.4 mm or more

(3) In the case of 980 MPa class TS: 980 MPa or more and less than 1180 MPa EL: 19% or more λ: 20% or more R: 3.0 mm or less Erichsen value: 10.0 mm or more

(4) In the case of 1180 MPa class TS: 1180 MPa or more and less than 1270 MPa EL: 15% or more λ: 20% or more R: 4.5 mm or less Eriksen value: 9.6 mm or more

  In the present invention, it is assumed that TS is 590 MPa or more and less than 1270 MPa. If TS is less than 590 MPa or 1270 MPa or more, even if EL, λ, R, and Erichsen values are good Treat as outside.

  The following Table 1 to Table 5 can be considered as follows. No. shown in Table 4 and Table 5 below. 1-4. Nos. 1, 3, 4, 11, 14, 15, 20, and 28 are examples of cooling with the above pattern i. Nos. 2 and 6 are examples cooled with the above pattern iii. Reference numeral 19 is an example of cooling with the pattern iv, and the rest is an example of cooling with the pattern ii.

  In Tables 4 and 5 below, examples in which a circle is attached to the overall evaluation are steel plates that satisfy the requirements defined in the present invention, and mechanical properties (EL, (λ, R, Erichsen values) are satisfied. Therefore, it can be seen that the high-strength steel sheet of the present invention is excellent in elongation and local deformability, and is good throughout the workability.

  On the other hand, the example (No. 4, 8, 9, 12, 15, 18-20, 31, 34-36) where x is attached to the comprehensive evaluation satisfies any of the requirements defined in the present invention. There is no steel plate. Details are as follows.

  No. No. 4 is an example in which the average cooling rate when cooling to an arbitrary temperature T satisfying the above formula (1) after being heated and held in the two-phase temperature range is too small, causing pearlite transformation, and the desired residual γ amount is Not obtained. Therefore, the strength is insufficient. No. No. 8 is an example in which the holding time in the two-phase temperature range is too short, and the amount of residual γ was not secured, so that the strength was insufficient. No. No. 9 is an example in which after soaking, the temperature is maintained at a temperature exceeding the temperature in the T1 temperature range defined in the present invention, not maintained in the T1 temperature range, but cooled to the T2 temperature range and maintained in this temperature range. Since a large amount of polygonal ferrite is generated, the amount of high-temperature region bainite generated is small and the amount of residual γ is also small. Therefore, the elongation and Erichsen values are lowered, and workability cannot be improved.

  No. No. 12 is an example in which, after being held in the T1 temperature range, cooled to a temperature lower than the T2 temperature range, it was not maintained in the T2 temperature range, and low temperature range bainite was hardly generated, and coarse MA mixing was observed by SEM observation. It was confirmed that a large amount of phase was present, and a lot of quenched martensite was present. Accordingly, all of the elongation, the hole expansion ratio, the critical bending radius, and the Erichsen value do not satisfy the acceptance criteria defined in the present invention, and the workability cannot be improved. No. 15 is an example in which the holding time in the T1 temperature range is long and not held in the T2 temperature range, and the generation of low temperature range bainite or the like is suppressed. Moreover, many coarse MA mixed phases are producing | generating. Accordingly, the hole expansion rate and the critical bending radius are small, the Erichsen value is also small, the local deformability is lowered, and the workability of the steel sheet cannot be improved.

  No. In No. 18, since the heating temperature is too high, polygonal ferrite is hardly produced, and the production amount of high temperature region bainite, low temperature region bainite and the like is excessive. Accordingly, the elongation is lowered and the workability of the steel sheet cannot be improved. No. No. 19 is an example in which the holding time in the T1 temperature range is too long and the cooling is performed without holding in the T2 temperature range, and the generation of low temperature range bainite or the like is suppressed. In addition, a large amount of coarse MA mixed phase was produced. Accordingly, the hole expansion rate is small, the Erichsen value is small, the local deformability is lowered, and the workability is deteriorated. No. No. 20 is an example in which after the soaking process, the temperature is not maintained in the T1 temperature range but is cooled to the T2 temperature range at a stretch, and is maintained at two temperatures in this temperature range. Since the temperature is maintained only in the T2 temperature range, almost no high-temperature range bainite is generated and almost no residual γ is generated. Accordingly, the elongation and the Erichsen value are lowered, and the workability is deteriorated. No. No. 31 was an example in which the heating temperature was too low, and the amount of polygonal ferrite produced increased, and high temperature region bainite, low temperature region bainite, etc., and residual γ were not produced at all. Accordingly, the elongation is lowered and the workability cannot be improved.

  No. No. 34 is an example in which the amount of C is too small. Since the amount of residual γ produced is too small, the elongation and Erichsen values are small, and the workability is deteriorated. No. No. 35 is an example in which the amount of Si is too small. Since the amount of residual γ produced is too small, elongation is lowered and workability is deteriorated. No. 36 is an example in which the amount of Mn is too small, and quenching has not been sufficiently performed. Thus, formation of polygonal ferrite is promoted during cooling, but generation of low-temperature region bainite and the like is suppressed. Accordingly, the elongation, the hole expansion rate, and the limit bending radius are reduced, and the workability is deteriorated.

  From the above results, it can be seen that according to the present invention, a high-strength steel sheet with improved workability can be provided.

  Next, among the 780 MPa class steel plates shown in Tables 4 and 5 above, examples satisfying the requirements defined in the present invention (Nos. 3, 5-7, 11, 14, 16, 17, 23, 23 26, 30, 32, 37 to 43), the relationship between tensile strength (TS) and elongation (EL) is shown in FIG. In FIG. 4, ● represents the result of the average equivalent circle diameter D of the polygonal ferrite grains being 10 μm or less, and ▪ represents the result of the average equivalent circle diameter D of the polygonal ferrite grains exceeding 10 μm.

  As is clear from FIG. 4, even if the tensile strength (TS) is the same, the elongation (EL) can be increased by suppressing the average equivalent circle diameter D of the polygonal ferrite grains to 10 μm or less, and the workability is further improved. I can see that

Claims (10)

% By mass
C: 0.10 to 0.3%,
Si: 1.0-3%
Mn: 1.0 to 2.5%
Al: 0.005 to 3% and P: 0.1% or less,
S: satisfying 0.05% or less,
The balance is a steel plate made of iron and inevitable impurities,
The metallographic structure of the steel sheet includes polygonal ferrite, bainite, tempered martensite, and retained austenite,
(1) When the metal structure is observed with a scanning electron microscope,
(1a) The area ratio a of the polygonal ferrite is more than 50% with respect to the entire metal structure, and (1b) the bainite is
High temperature zone bainite having an average interval between adjacent retained austenite and / or carbide of 1 μm or more;
It is composed of a composite structure with low-temperature region-generated bainite in which the average interval between adjacent retained austenite and / or carbide is less than 1 μm,
The area ratio b of the high temperature region bainite is 5 to 40% with respect to the entire metal structure,
The total area ratio c of the low temperature region bainite and the tempered martensite satisfies 5 to 40% with respect to the entire metal structure,
(2) A high-strength steel sheet excellent in workability, wherein the volume ratio of the retained austenite measured by a saturation magnetization method is 5% or more with respect to the entire metal structure. When an MA mixed phase in which quenched martensite and retained austenite are present when the metallographic structure is observed with an optical microscope, the equivalent circle diameter in the observation cross section with respect to the total number of MA mixed phases. 2. The high-strength steel sheet according to claim 1, wherein the number ratio of MA mixed phases satisfying d of greater than 7 μm is less than 15% (including 0%). 3. The high-strength steel sheet according to claim 1, wherein an average equivalent circle diameter D of the polygonal ferrite grains is 10 μm or less (not including 0 μm). The steel sheet, as another element,
The high-strength steel sheet according to any one of claims 1 to 3, comprising Cr: 1% or less (not including 0%) and / or Mo: 1% or less (not including 0%). The steel sheet, as another element,
Ti: 0.15% or less (excluding 0%),
5. The composition according to claim 1, comprising at least one element selected from the group consisting of Nb: 0.15% or less (excluding 0%) and V: 0.15% or less (excluding 0%). The high strength steel plate according to any one of the above. The steel sheet, as another element,
The high-strength steel sheet according to any one of claims 1 to 5, containing Cu: 1% or less (not including 0%) and / or Ni: 1% or less (not including 0%). The steel sheet, as another element,
B: The high-strength steel plate according to any one of claims 1 to 6, containing 0.005% or less (excluding 0%). The steel sheet, as another element,
Ca: 0.01% or less (excluding 0%),
8. One or more elements selected from the group consisting of Mg: 0.01% or less (not including 0%) and rare earth elements: 0.01% or less (not including 0%) A high-strength steel sheet according to any one of the above. The high-strength steel plate according to any one of claims 1 to 8, which has a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel plate. A method for producing the high-strength steel sheet according to any one of claims 1 to 8,
A step of heating to a temperature range of Ac ₁ point + 20 ° C. or higher and Ac ₃ point + 20 ° C. or lower;
Holding for 50 seconds or more in the temperature range;
A step of cooling at an average cooling rate of 2 to 50 ° C./second to an arbitrary temperature T satisfying the following formula (1), a step of holding for 10 to 100 seconds in a temperature range satisfying the following formula (1),
Holding for 200 seconds or more in a temperature range satisfying the following formula (2);
In this order, a method for producing a high-strength steel sheet excellent in workability.
400 ° C. ≦ T1 (° C.) ≦ 540 ° C. (1)
200 ° C. ≦ T2 (° C.) <400 ° C. (2)