JP5644095B2 - High strength steel sheet having good tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance, manufacturing method of high strength cold rolled steel sheet, manufacturing method of high strength galvanized steel sheet - Google Patents

Description

  The present invention relates to a high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance, a method for producing a high-strength cold-rolled steel sheet, and a method for producing a high-strength galvanized steel sheet.

  In recent years, there has been an increasing demand for higher strength steel sheets used in automobiles and buildings. High-strength cold-rolled steel sheets having a maximum tensile strength of 900 MPa or more have been rapidly applied mainly to reinforcing materials such as bumpers and impact beams. A problem in using ultra-high strength steel is a phenomenon called delayed fracture. Delayed fracture is caused by stress applied to the steel material or hydrogen embrittlement, and is a phenomenon in which hydrogen is diffused into the stress concentration part of the steel material used as the structure, thereby causing the structure to be destroyed. Specifically, for example, a phenomenon such as a PC steel wire or a bolt that suddenly breaks a member on which high stress acts under a use situation.

  Delayed destruction is known to be closely related to hydrogen entering from the environment. There are various types of hydrogen that enter the steel from the environment, such as hydrogen contained in the atmosphere and hydrogen generated in a corrosive environment. In any case, when hydrogen penetrates into the steel material, it can cause delayed fracture. From this, it is desirable to use steel in an environment where hydrogen is not present. However, in consideration of application to structures or automobiles, it will be used outdoors. Unavoidable.

  In addition, stress acting on the steel material used as the structure includes a stress applied to the structure and a residual stress in which a part of the stress generated when the structure is formed remains. In particular, in a structure used as a member after forming, such as a thin steel plate for automobiles, the residual stress becomes a big problem as compared with a thick plate or a steel bar that uses a product such as a bolt or a thick plate as it is. Therefore, when forming a steel sheet in which delayed fracture is a problem, a forming method that does not leave residual stress is desired.

  For example, Patent Document 1 describes a technique in which a steel sheet is once heated and processed to a high temperature and then baked using a mold to increase the strength. In this method, steel is processed at a high temperature, so dislocations introduced during processing that cause residual stress are recovered, or transformation occurs after processing and the residual stress is relaxed. It is known that not much remains. Therefore, as described above, the delayed fracture characteristics can be improved by performing hot processing and strengthening the steel sheet using the subsequent quenching.

  Further, in machining such as cutting and punching, there is a concern of causing delayed fracture because residual stress exists on the cut surface. For this reason, at the time of processing a high-strength steel plate of 980 MPa or more, residual stress generation is avoided by using a technique that does not involve machining such as laser for cutting. However, there is a problem that laser cutting is expensive compared to shear cutting and punching.

  In response to these problems, steel materials capable of avoiding delayed fracture have been developed by improving the hydrogen embrittlement resistance of steel materials. For example, Non-Patent Document 1 shows temper softening resistance such as Cr, Mo and V by quenching a steel material from a high-temperature austenite single phase to form a martensite single phase structure and then performing a tempering treatment. A high-strength bolt in which fine precipitates of additive elements are finely precipitated in martensite in a consistent manner to improve the hydrogen embrittlement resistance of steel is described. In this method, hydrogen that has penetrated into the steel material is trapped around VC or the like that has been deposited in the martensite in a consistent manner. It is suppressed from diffusing or concentrating on the starting point. Utilizing this technique, the development of steel sheets with high strength and excellent delayed fracture resistance has been underway.

  Improvement in delayed fracture resistance utilizing hydrogen trap sites such as VC is brought about by the consistent deposition of these precipitates in the martensite structure. Therefore, it is essential to make these precipitates consistently precipitate in the structure. However, precipitation of these precipitates requires precipitation heat treatment for several hours or more, and there is a problem in manufacturability. That is, in steel sheets manufactured using general thin steel sheet manufacturing equipment such as continuous annealing equipment and continuous hot dip galvanizing equipment, the structure is controlled in a short time of at most several tens of minutes. There is a problem that it is difficult to obtain an effect of improving delayed fracture resistance due to an object.

  In addition, when utilizing precipitates precipitated in the hot rolling process, even if these precipitates are precipitated in the hot rolling process, the steel sheet is processed during subsequent cold rolling and recrystallized during continuous annealing. Thus, the orientation relationship between the precipitate and the ferrite and martensite as the parent phase is lost, that is, the precipitate is not matched. As a result, the delayed fracture resistance of the obtained steel sheet is also greatly reduced.

  Further, the steel sheet structure of a high-strength steel sheet that is likely to cause delayed fracture is a structure mainly composed of martensite. Since the temperature at which the martensite structure is formed is low, precipitates that serve as hydrogen trap sites such as VC cannot be deposited in the temperature range where the martensite structure is formed. As a result, when it is intended to improve the delayed fracture resistance due to hydrogen traps due to matched precipitates such as VC in thin steel sheets, the steel structure is once built by continuous annealing equipment or continuous hot dip galvanizing equipment, and then additional It is necessary to perform heat treatment to precipitate these precipitates, resulting in a significant increase in manufacturing cost. In addition, when an additional heat treatment is performed on a structure mainly composed of martensite, a problem of softening occurs. As a result, it is difficult to utilize matched precipitates for improving delayed fracture resistance of high strength thin steel sheets.

  Moreover, since the steel described in Non-Patent Document 1 has a C content of 0.4% or more and contains a large amount of alloying elements, the workability and weldability required for thin steel sheets are poor.

  Patent Document 2 describes a technique for preventing hydrogen defects by an oxide mainly composed of Ti and Mg. However, although the technique described in Patent Document 1 is a thick steel plate, and consideration is given to hydrogen embrittlement after welding with high heat input, both high formability and hydrogen embrittlement resistance required for a thin steel plate are achieved. Is not considered at all.

Conventionally, a thin steel plate is (1) a thin plate, so even if hydrogen enters, it is released in a short time, and (2) there is almost no use of a steel plate of 900 MPa or more in terms of workability. The problem with brittleness was small. However, since the demand for the application of high-strength steel plates is rapidly increasing, it is necessary to develop high-strength steel plates with excellent hydrogen embrittlement resistance.
However, techniques for improving hydrogen embrittlement resistance have been developed for steel materials that are often used as products such as bolts, strips, and thick plates, and are often used below the yield strength or yield stress. That is, it was not a technology that considered steel materials that required hydrogen embrittlement resistance as well as workability such as cutting and member molding (press molding) like automobile members. In particular, a stress called residual stress remains in the member after molding. Although the residual stress is local, it may be high enough to exceed the yield stress of the material. For this reason, it is required that hydrogen embrittlement does not occur under high residual stress.

  Regarding the hydrogen embrittlement of thin steel sheets, for example, Non-Patent Document 2 reports the promotion of hydrogen embrittlement due to the processing-induced transformation of the amount of retained austenite. This is in consideration of the forming process of a thin steel sheet, but the regulation of the amount of retained austenite that does not deteriorate the hydrogen embrittlement resistance is described. That is, it relates to a high-strength thin steel sheet having a specific structure, and is not a fundamental measure for improving hydrogen embrittlement resistance.

  In general, as a technique for ensuring a high strength of 900 MPa or more, organization strengthening utilizing martensite is known. Examples of techniques for increasing the strength of martensite include techniques described in Non-Patent Documents 3 to 5. Non-Patent Document 3 discloses that the hardness (strength) of a lath-like structure (lass martensite structure), which is one of the martensite structure shapes, is the amount of solute C in martensite, crystal grain size, and carbide. It is described that it depends on precipitation strengthening and dislocation strengthening. Non-Patent Document 4 and Non-Patent Document 5 describe that the strength of martensite can be greatly increased by refining the crystal grain size, particularly the block size, which is one of the structural units constituting martensite. Yes.

  Moreover, as a prior art regarding a high strength steel plate, the technique of patent document 3-patent document 9 is mentioned, for example. Moreover, as a conventional technique regarding the hot dip galvanized steel sheet, for example, a technique described in Patent Document 10 can be cited.

Japanese Patent Laid-Open No. 2002-18531 JP-A-11-293383 Japanese Patent Laid-Open No. 11-100638 Japanese Patent Application Laid-Open No. 11-296991 Japanese Patent Laid-Open No. 9-13147 JP 2002-363695 A JP 2003-105514 A JP 2003-213369 A JP 2003-213370 A JP 2002-97560 A

New development of hydrogen embrittlement elucidation, Japan Iron and Steel Institute, published in January 1997 CAMP-ISIJ, vol. 5, no. 6, pp. 1839-1842, Yamazaki et al., October 1992, published by Japan Iron and Steel Institute F. B. Pickering: Hardenability Concepts with Applications to Steel AIME (1978), p179 M.M. Wang: IS3-2007, p203 T. T. et al. Ohmura: IS3-2007, p35

  As a technique for improving the delayed fracture resistance of the molded member, it is conceivable to use a molding technique for reducing the residual stress of the molded member. However, since the residual stress remains in the molded member, there are many portions depending on the characteristics of the steel material, and the member shape cannot be changed greatly. For this reason, there are many problems in the technique of reducing the residual stress of the molded member using the molding method.

  Further, in order to improve the residual stress and hydrogen embrittlement resistance after forming, it is conceivable to use a method for controlling the characteristics of the steel material. That is, by lowering the yield stress of the steel sheet and increasing the uniform elongation of the steel sheet, the yield strength when the same plastic deformation is applied is kept low, and the residual stress of the member after forming is reduced. The member is considered to have excellent hydrogen embrittlement resistance. For this reason, in order to increase the hydrogen embrittlement resistance after forming a high-strength steel sheet, it is desirable that the yield stress is low and the uniform elongation is large.

  However, a steel sheet that has secured a high strength of 900 MPa or more by utilizing a martensite structure has a high strength but has a very low uniform elongation and poor workability. There is stress. Since the residual stress is a part of the stress at the time of forming remaining in the steel material, the higher the strength of the steel plate, the higher the strength of the steel sheet that becomes high stress at the time of forming, the higher the residual stress. For this reason, the high-strength steel sheet using the martensite structure has a problem that the hydrogen embrittlement resistance of the formed member is insufficient and the risk of delayed fracture is high.

As a technique for solving this problem, it is conceivable to use a composite structure steel plate having a main phase of ferrite and a hard structure of martensite. In such a composite structure steel plate, ductility is ensured with soft ferrite and strength is ensured with hard martensite. However, since the main phase is soft ferrite, the yield stress can be greatly reduced.
However, in a multiphase steel sheet composed of ferrite and martensite, in order to ensure a tensile strength of 900 MPa or more of the steel sheet, it is necessary to sufficiently increase the volume ratio of martensite, which causes deterioration of uniform elongation and an increase in yield stress. It was inevitable.

  In order to solve this problem, it is conceivable to improve the strength of the martensite structure to secure a strength of 900 MPa or more even with a small amount of martensite. Since the strength of the martensite structure strongly depends on the steel plate component (particularly C), it is difficult to increase the strength without changing the steel plate component. However, when the addition amount of C is increased, weldability including spot welding is deteriorated. For this reason, in the case of a thin steel plate using spot or laser welding as a fastening means for a formed member, or a steel plate that is formed after welding, such as a tailored blank material, the amount of C added is suppressed to ensure sufficient weldability. I had to. Therefore, when producing a high strength steel sheet with a maximum tensile strength of 900 MPa class, C is added in an addition amount suppressed to a level where sufficient weldability is obtained, and the volume ratio of martensite is increased to increase the strength. Had to secure. As a result, in conventional high strength steel plates with a maximum tensile strength of 900 MPa, deterioration of uniform elongation and increase in yield stress were inevitable.

  In addition, when performing processes such as continuous annealing and continuous hot dip galvanizing lines, inevitable heat treatment such as annealing, tempering in the overaging zone after cooling, and alloying heat treatment after immersion in the plating bath is additionally provided. By doing so, the martensitic structure is tempered and softened, so that it is difficult to ensure the strength of the steel sheet. From this, when manufacturing a high-strength steel sheet exceeding 900 MPa, for example, using a manufacturing facility that receives such additional heat treatment, generally, even if softening occurs due to tempering of the martensite structure. In order to ensure sufficient strength, the martensite volume fraction is further increased. This results in a further increase in yield stress.

  As described above, in the conventional technique, the strength is ensured by increasing the volume ratio of martensite. However, when the volume ratio of martensite is increased, the yield stress (YS) increases, so the yield ratio (YR), which is the ratio of the maximum tensile strength (TS) to the yield stress (YS), increases, and delay resistance Destructive properties deteriorate. In addition, when the volume ratio of martensite is increased, the volume ratio of ferrite that has contributed to ensuring ductility is relatively decreased, and it is inevitable that ductility is greatly deteriorated and workability is lowered. .

  The present invention has been made in view of such circumstances, and without increasing the volume ratio of martensite, which is a hard structure, by ensuring a high strength of a tensile maximum strength of 900 MPa or more, the yield ratio ( It is an object of the present invention to provide a high-strength steel sheet having a small YR), sufficient delayed fracture resistance, sufficient ductility and high strength, and a method for producing a high-strength steel sheet.

  The present inventor confronted the conventional common sense of enjoying high strength in exchange for increasing ductility deterioration and yield ratio (YR) by increasing the volume ratio of martensite, and yield ratio (YR) is In order to realize a high-strength steel sheet that is small, has sufficient ductility, and can obtain high strength, intensive studies have been made. As a result, by increasing the strength of the martensite structure without increasing the volume ratio of martensite, which is a hard structure, it is possible to minimize the decrease in ferrite volume that contributes to securing ductility, and sufficient ductility, The inventors have found a method capable of ensuring sufficient delayed fracture resistance and securing a high strength with a maximum tensile strength of 900 MPa or more, and have conceived the high-strength cold-rolled steel sheet and the high-strength galvanized steel sheet of the present invention.

The gist of the present invention is as follows.
(1) By mass%, C: 0.07 to 0.25%, Si: 0.3 to 2.50%, Mn: 1.5 to 3.0%, Ti: 0.005 to 0.09% , B: 0.0001-0.01%, P: 0.001-0.03%, S: 0.0001-0.01%, Al: 2.5% or less, N: 0.0005-0. 0100%, O: 0.0005-0.007%, the balance is steel consisting of iron and inevitable impurities,
The steel sheet structure is mainly composed of ferrite, contains martensite , the ferrite volume fraction is 50% or more, the martensite block size is 1 μm or less, and the C concentration in the martensite is 0.3% to 0.9%. %, Yield ratio (YR) consisting of ratio of maximum tensile strength (TS) and yield stress (YS) is 0.75 or less, good tensile maximum strength with good ductility and delayed fracture resistance A high-strength steel sheet having 900 MPa or more.

(2) Further, Nb: 0.005 to 0.09% by mass%, high tensile strength of 900 MPa or more with good ductility and delayed fracture resistance according to (1) Strength steel plate.
(3) Further, by mass%, Cr: 0.01-2.0%, Ni: 0.01-2.0%, Cu: 0.01-2.0%, Mo: 0.01-0. The high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to (1) or (2), comprising 8% of one kind or two or more kinds.
(4) Furthermore, it is contained in mass%, and V: 0.005 to 0.09%. The ductility and delayed fracture resistance described in any one of (1) to (3) are good. A high-strength steel sheet having a maximum tensile strength of 900 MPa or more.
(5) Any one of (1) to (4), further comprising 0.0001 to 0.5% in total of one or more of Ca, Ce, Mg, and REM in mass% A high-strength steel sheet having good tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to item 1.

(6) A high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to any one of (1) to (5), wherein the surface has a galvanized layer.

  In addition, the inventor has intensively studied a method for producing a high-strength steel sheet that can increase the strength of the martensite structure without increasing the volume ratio of the martensite that is a hard structure. As a result, in the annealing process after cold rolling, fine precipitates such as Ti, Nb, and V are deposited to suppress the growth of martensite laths forming individual blocks, and to the steel sheet after annealing. By introducing dislocations as nucleation sites, laths having a plurality of different orientations are generated, so that the martensite block size is 1 μm or less and the tensile maximum strength is 900 MPa or more with good ductility and delayed fracture resistance. The inventors have found that a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet according to the present invention can be obtained, and have conceived the production method of the present invention.

(7) (1) to (5) be any process for producing a high strength steel sheet according to the, casting a slab having the chemical components described in any of (1) to (5), directly or Once cooled, heated to 1050 ° C. or higher, completed hot rolling above the Ar ₃ transformation point, wound up in a temperature range of 400 ° C. to 670 ° C., pickled, and then cold-rolled with a rolling reduction of 40-70% In order to pass through the continuous annealing line,
After heating between 550 ° C. and 760 ° C. for 30 seconds or longer and annealing at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., cooling between a maximum heating temperature of 630 ° C. and an average cooling rate of 14 ° C./sec or less is performed. , The temperature between 630 ° C. and 570 ° C. is cooled at an average cooling rate of 3 ° C./second or more, and bending-bending deformation with an indentation amount of 1 mm or more is performed using a roll having a roll diameter of 350 mm or less in a temperature range of 700 ° C. to 400 ° C. A method for producing a high-strength cold-rolled steel sheet having a maximum tensile strength of 900 MPa or more with good ductility and delayed fracture resistance.

(8) The ductility according to (7), characterized in that after cooling between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more, the temperature is maintained at 450 ° C. to 250 ° C. for 30 seconds or more. And a method for producing a high-strength cold-rolled steel sheet having a tensile maximum strength of 900 MPa or more with good delayed fracture resistance.
(9) The inside of the annealing furnace of the continuous annealing line contains ₁ to 60% by volume of H ₂ , and is an atmosphere composed of the balance N ₂ , H ₂ O, O ₂ and unavoidable impurities, and the moisture pressure in the atmosphere The logarithm log (PH ₂ O / PH ₂ ) of the hydrogen partial pressure and -3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5 is described in (7) or (8) A method for producing a high-strength cold-rolled steel sheet having a maximum tensile strength of 900 MPa or more with good ductility and delayed fracture resistance.
(10) After producing a high-strength cold-rolled steel sheet by the method for producing a high-strength cold-rolled steel sheet according to any one of (7) to (9), zinc electroplating is performed and ductility and resistance A method for producing a high-strength galvanized steel sheet having a maximum tensile strength of 900 MPa or more with good delayed fracture characteristics.

(11) A method for producing a high-strength steel sheet according to (6), wherein the slab having the chemical component according to any one of (1) to (5) is cast and directly or once cooled, and then 1050 ° C. or higher. To complete the hot rolling at the Ar ₃ transformation point or higher, winding in a temperature range of 400 ° C. to 670 ° C., pickling, cold rolling with a rolling reduction of 40 to 70%, continuous hot dip galvanization When letting through the line,
After heating at 550 ° C. to 760 ° C. for 30 seconds or more during heating and annealing at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., cooling between the maximum heating temperature to 630 ° C. at an average cooling rate of 14 ° C./second or less is performed. ℃ ~ [(Zinc plating bath temperature -40 ℃) ~ (Zinc plating bath temperature + 50 ℃)] is cooled at an average cooling rate of 3 ℃ / second or more, roll diameter 350mm in the temperature range of 700 ℃ ~ 400 ℃ Using the following rolls, a bending-unbending deformation with an indentation amount of 1 mm or more is performed, followed by immersion in a galvanizing bath and cooling. A tensile strength of 900 MPa or more with good ductility and delayed fracture resistance A method for producing a high-strength galvanized steel sheet having

(12) After dipping in the galvanizing bath, alloying is performed at a temperature of 460 ° C. to 600 ° C., and cooling is performed. Tensile strength with good ductility and delayed fracture resistance according to (11) A method for producing a high-strength galvanized steel sheet having a maximum strength of 900 MPa or more.
(13) the said annealing furnace of a continuous galvanizing line, and H ₂ contains 1 to 60 vol%, the remainder _{N 2,} H 2 O, an atmosphere consisting of _{O 2} and inevitable impurities, in the atmosphere (11) or (12) characterized in that the logarithm log of water pressure and hydrogen partial pressure (PH ₂ O / PH ₂ ) is −3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5 A method for producing a high-strength galvanized steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to any one of the items.

  According to the present invention, it is possible to ensure a high tensile strength of 900 MPa or higher without increasing the volume fraction of martensite, which is a hard structure, a low yield ratio (YR), and sufficient delayed fracture resistance. It is possible to provide a high-strength cold-rolled steel sheet and a high-strength galvanized steel sheet having sufficient ductility and high strength and a method for producing the same.

It is a stress-strain curve. FIG. 2 is a view for explaining the shape of the test piece after being flange-formed.

The high-strength steel sheet of the present invention has good ductility and delayed fracture resistance and has a maximum tensile strength of 900 MPa or more.
The steel sheet structure of the high-strength steel sheet of the present invention is mainly composed of ferrite and contains martensite, and may contain bainite. In the steel sheet structure of the high-strength steel sheet of the present invention, the volume ratio of ferrite is 50% or more. If the volume fraction of ferrite is less than 50%, the volume fraction of martensite is relatively high, the yield stress (YS) increases, the yield ratio (YR) increases, and the residual stress increases. The delayed fracture resistance cannot be sufficiently obtained.

In the steel sheet structure of the high-strength steel sheet of the present invention, the ferrite volume fraction needs to be as high as possible from the viewpoint of ensuring ductility. However, even if the martensite is hardened, there is a limit to the strength that can be obtained, so the required volume fraction of ferrite varies depending on the maximum tensile strength.
For example, in the case of a high-strength steel plate having a maximum tensile strength of 900 to 1130 MPa, the ferrite volume fraction is preferably in the range of 60% to 85%, and more preferably in the range of 65% to 80%.
In the case of a high-strength steel sheet having a maximum tensile strength of 1130 to 1280 MPa, the ferrite volume fraction is preferably in the range of 55% to 80%, and more preferably in the range of 60% to 75%.
In the case of a high-strength steel sheet having a maximum tensile strength of 1280 to 1580 MPa, the volume ratio of ferrite is preferably in the range of 50% to 75%, and more preferably in the range of 55% to 70%.
By controlling the volume ratio of the ferrite within the above range, a strength-ductility balance (TS ×) which is a product of high strength with a maximum tensile strength of 900 MPa or more, maximum tensile strength (TS) and total elongation (El.) In a tensile test. El.) Can be 18000 (MPa ×%) or more, and excellent ductility can be provided at the same time.

  On the other hand, for example, in conventional steel, the strength is increased by increasing the volume ratio of martensite. Therefore, in steel having a maximum tensile strength of about 590 MPa, the volume ratio of ferrite is about 90%. In the steel having the maximum strength of about 980 MPa, the volume fraction of ferrite is about 50%, and in the steel having the maximum tensile strength of about 1180 MPa, the ferrite is hardly contained.

In the present invention, the average crystal grain size (dF) of ferrite is preferably 5 μm or less, more preferably 4.5 μm or less, and most preferably 4 μm or less.
The reason why the average crystal grain size (dF) of the ferrite is set in the above range is that by reducing the crystal grain size, the maximum tensile strength is increased without significantly reducing the ductility. If the average crystal grain size of ferrite exceeds the above range, the contribution of strength increase due to fine graining becomes small, so the strength must be compensated for by increasing the volume ratio of martensite, and the volume ratio of ferrite decreases. And this causes undesirable ductile deterioration. In addition, when the average crystal grain size of ferrite is 5 μm or less, localization of deformation and crack propagation are less likely to occur. If tensile deformation, local ductility is improved. If bending or hole expansion molding is used, bendability is increased. And the hole expansion rate is improved. However, since the decrease in ferrite grain size leads to an increase in yield stress, it is not preferable to extremely reduce the average crystal grain size of ferrite, and the lower limit is preferably 1 μm or more.

The average crystal grain size (dM) of martensite is preferably 1/3 or less of the average crystal grain size (dF) of ferrite, specifically 1.5 μm or less. More preferably, it is set to 2 μm or less, and most preferably 0.9 μm or less.
The reason why the average crystal grain size (dM) of martensite is in the above range is to suppress the concentration of deformation at the interface between ferrite and martensite during deformation and to suppress the formation of microvoids and cracks at the interface. It is. That is, soft ferrite and hard martensite are greatly different in deformability, so if they are tensile deformation under large deformation after necking, if bending molding, during bending with small R, or during hole expansion molding Furthermore, deformation concentrates on the interface between ferrite and martensite, leading to destruction. Therefore, these characteristics can be greatly improved by making the martensite fine and suppressing the concentration of deformation at each interface. In addition, when machining such as shear cutting or punching, if coarse martensite is present in the cut part, flaws and fine cracks originating from martensite are generated, which is subsequently performed. This greatly reduces workability during processing. For this reason, it is preferable to make the martensite particle size as small as possible.

  In the high-strength steel sheet of the present invention, the C concentration in martensite is 0.3% to 0.9%, preferably 0.35% to 0.8%, It is more preferable that it is% to 0.7%. If the C concentration in the martensite is less than the above range, the effect of strengthening the martensite due to C in the martensite cannot be sufficiently obtained, and the strength of the high-strength steel sheet is insufficient. Moreover, when the C concentration in martensite exceeds the above range, weldability and workability become insufficient.

  The C concentration in martensite is determined by, for example, the following method. First, a steel plate is embedded and polished in a cross section parallel to the rolling direction. Then, the structure | tissue observation by SEM is performed, the position of each structure | tissue is specified, and C density | concentration in a martensite is measured by EPMA (Electron Probe Micro Analyzer). Although the C concentration in martensite can be measured by point analysis, it is preferable to measure by obtaining a C concentration distribution in the tissue by performing surface analysis. Generally, ferrite contains almost no C, but martensite contains a large amount of C. From this, if the optical microscope confirms that the steel sheet structure is composed of ferrite and martensite in advance, it is possible to measure the C concentration in martensite even if only component mapping by EPMA is used. It is. However, since the high strength steel sheet of the present invention controls the precipitation of Ti carbide, the crystal grain sizes of ferrite and martensite are small. Therefore, it is desirable that the step size when performing surface analysis is as small as possible, and it is desirable that the step size be 1 μm or less. However, if the step size is excessively small, it takes a long time to measure, so it is desirable that the step size be 0.05 μm or more.

  Further, the martensite contained in the steel sheet structure of the high-strength steel sheet of the present invention is granular and can be present either in the ferrite grain boundaries or in the grains. The martensite organization is composed of a hierarchical structure called packets, blocks, and laths.

  The packet is composed of a plurality of blocks having the same crystal habit plane of parent phase austenite and different crystal orientations. One packet is composed of a maximum of six blocks having a KS relationship (Kurdjumov-Sachs relationship). Note that in general optical microscope observation, blocks having variants with small crystal orientation differences (brother crystals with the same crystal structure but different crystal orientations) cannot be distinguished, so two variant pairs with crystal orientation differences of 11 ° or less are present. It is defined as one block. In this case, one packet is composed of three blocks.

A block is a collection of laths having the same crystal orientation. Although the boundaries of the individual laths exist inside the block, the laths within the individual blocks have the same crystal orientation, so that they work as if they were one grain when deformed. Therefore, the strength of martensite depends on the block size.
The lath is needle-shaped and granular depending on the observation direction. Further, the lath has a KS relationship with the austenite of the parent phase, and there are those having 24 orientation relationships with the parent phase.

  Packets, blocks, and laths that make up the martensitic structure can be observed with a transmission electron microscope (TEM), crystal orientation analysis using a scanning electron microscope (SEM) -backscattered electron diffraction (EBSP) method, field emission scanning This can be confirmed by high resolution crystal orientation analysis using an electron microscope (FE (Field Emission) -SEM) -EBSP method.

The present inventor has focused on the relationship between the martensite block size present in the steel sheet and the strength of the steel sheet, and is utilized as a reinforcing structure of the steel sheet by making the martensite block size finer. It has been found that the strength of martensite can be further increased, and it has been found that it is possible to increase the strength exceeding 900 MPa while securing the volume fraction of ferrite capable of ensuring a low yield stress. Specifically, by setting the martensite block size to 1 μm or less, it is possible to utilize martensite as a strengthened structure more than ever.
The martensite constituting the steel sheet structure of the high-strength steel sheet of the present invention is composed of a block size of 1 μm or less. When the block size exceeds 1 μm, the effect of strengthening martensite is not sufficiently obtained, and the strength of the high-strength steel sheet is insufficient.

  For example, in the conventional steel, the block size of martensite having the same crystal orientation was extremely large, from several μm to several tens of μm. For this reason, in the conventional steel, for example, when the steel sheet structure is controlled to be several μm or less, the size of individual martensite grains utilized as the reinforcing structure of the steel sheet is several μm or less, and each martensite grain is single. It consists of blocks. As a result, in conventional steel, martensite could not be fully utilized as a strengthened structure.

  The martensite block size can be measured by TEM observation, SEM-EBSP method, FE-SEM-EBSP method, or the like. The SEM-EBSP method is preferable because a wide area can be measured in a short time. In addition, the FE-SEM-EBSP method is desirable because it can perform tissue observation with high accuracy.

When tissue observation is performed using the FE-SEM-EBSP method, the interval between the measurement points may be in a range in which tissue orientation analysis is possible, but is preferably 10 nm to 200 nm. If the interval between the measurement points exceeds 200 nm, the accuracy may decrease, and an accurate block size may not be measured. Further, if the interval between the measurement points is less than 10 nm, the measurement time is extremely long although it is desirable from the point of accuracy.
In the present invention, when measuring the martensite block size using the FE-SEM-EBSP method, it is possible to analyze the orientation of the structure by measuring the block size of several martensite grains in a preliminary experiment. It is preferable to carry out after determining the interval between the measurement points.

The high strength steel sheet of the present invention has a yield ratio (YR) composed of a ratio of maximum tensile strength (TS) and yield stress (YS) of 0.75 or less, preferably in the range of 0.5 to 0.725. Yes, more preferably in the range of 0.55 to 0.7.
The reason why the yield ratio is set to 0.75 or less is to reduce the stress acting on each part during molding. That is, the residual stress is generated when some of the dislocations introduced during forming remain in the steel sheet after forming. Therefore, residual stress can be reduced by reducing the stress acting on each part during molding. If the yield ratio (YR) exceeds 0.75, the yield strength after molding increases and the residual stress increases, so that the delayed fracture resistance cannot be sufficiently obtained.
Note that the amount of strain introduced into each part of the member during molding is determined by the member shape. Therefore, if the member shape is the same, the same strain is introduced regardless of the strength of the steel sheet. Moreover, since a high strength steel plate is often used in a relatively simple shape, the amount of strain applied to a formed member is usually small. For this reason, the residual stress when compared with the same strain amount becomes lower as the steel has a lower yield ratio (YR), and the steel having a lower yield ratio (YR) becomes more excellent in delayed fracture characteristics.

  The high-strength steel sheet of the present invention is mainly composed of ferrite in order to lower the yield ratio and contains martensite in order to ensure high strength. Therefore, the high-strength steel sheet has high strength and extremely high uniform elongation. In the present invention, it is desirable that the uniform elongation of the high-strength steel sheet is as large as possible as shown below. That is, a steel sheet with a small uniform elongation and a low yield ratio will be greatly strengthened with a small amount of deformation when applied to a site with large deformation, so the stress during forming will be close to the maximum tensile stress. . Therefore, with a steel sheet having a small uniform elongation, even if the yield ratio is lowered, the residual stress cannot be reduced and the effect of improving delayed fracture characteristics due to this cannot be obtained. Therefore, in order to improve the delayed fracture resistance of the high strength steel sheet, it is desirable that the uniform elongation is as large as possible. Specifically, when considering application of high-strength steel sheets to automotive structural members, the uniform elongation is preferably more than 5%, and more preferably 7% or more.

In the high-strength steel sheet of the present invention, delayed fracture characteristics can be evaluated by, for example, the following method.
First, a rectangular test piece obtained by cutting a high-strength steel plate in a direction perpendicular to the rolling direction is prepared, and the residual stress introduced at the time of cutting is removed by mechanical grinding of the end of the test piece. This is to remove the residual stress introduced at the time of cutting and the burrs and unevenness of the end face caused by the cutting, and to accurately evaluate the relationship between the residual stress introduced by the processing and the characteristics of the steel sheet and the delayed fracture characteristics. . If mechanical grinding is not performed, the delayed fracture characteristics evaluation results may vary due to uneven burr and end face irregularities due to shear wear, and accurate evaluation may not be possible. For this reason, the end of the test piece needs to be mechanically ground. After that, bending is performed at the center of the test piece in the length direction, the cross section in the width direction is made substantially L-shaped, and the end in the width direction of the test piece is flanged up so that the cross section in the length direction is made substantially U-shaped. A type.
Thereafter, the test piece is immersed in 0.5 mol / l sulfuric acid and electrolyzed with an electric current for 2 hours to allow hydrogen to penetrate into the test piece. Then, it can carry out by the method of taking out a test piece from a solution and evaluating generation | occurrence | production of the crack of a test piece visually.

In addition, the shape of the test piece before hydrogen penetrate | invades into a test piece should just be a shape which can evaluate the delayed fracture characteristic after shaping | molding, and is not specifically limited. The shape of the test piece before allowing hydrogen to penetrate into the test piece is, for example, the shape of a cylindrical drawing member, the shape of a member after a bending test, the shape of a member after an overhang test, or a combination of these. It does not matter.
In addition, delayed fracture characteristics are affected by residual stresses and defects caused by machining such as shear cutting, so before evaluating the occurrence of cracks in the test piece, it is necessary to perform mechanical grinding on the plate end of the test piece. Although it is preferable to perform machining, it is not necessary to perform machining on the plate end portion of the test piece as long as the delayed fracture characteristics can be evaluated.

Next, the chemical component (composition) of the high-strength steel sheet of the present invention will be described. In the following description, [%] is [% by mass].
“C: 0.07 to 0.25%”
C increases the strength of martensite and is included to increase the strength of the high-strength steel plate. However, when the C content exceeds 0.25%, weldability and workability become insufficient. On the other hand, if the C content is less than 0.07%, the strength is insufficient. The C content is preferably in the range of 0.08 to 0.23%, and more preferably in the range of 0.09 to 0.21%.

“Si: 0.3-2.50%” “Al: 2.5% or less”
Si and Al are ferrite stabilizing elements and increase the Ac ₃ transformation point, so that a large amount of ferrite can be formed at a wide annealing temperature, and are contained from the viewpoint of improving the structure controllability. Moreover, since Si and Al also contribute to solid solution strengthening, it is desirable to contain them positively. Since ferrite is a bcc phase containing almost no C, it is possible to concentrate C in austenite by forming a large amount of ferrite. Concentrating C in austenite contributes to increasing the strength of martensite.

If the Si content is less than 0.3%, the volume fraction of ferrite becomes insufficient, and the ductility, bendability and strength of the high-strength steel sheet become insufficient. When Al is contained, the same effect as that obtained when Si is contained can be obtained. However, when the above effect can be sufficiently obtained by containing only Si, Al may not be contained. . Further, if the Si content exceeds 2.50% or the Al content exceeds 2.5%, weldability and workability become insufficient.
The Si content is preferably in the range of 0.45 to 2.35%, and more preferably in the range of 0.6 to 2.2%. The content of Al is preferably in the range of 0.005 to 1.6%, and more preferably in the range of 0.01 to 0.6%. Moreover, since Al can be utilized as a deoxidizing material, it may be contained not only for the structure control of the steel sheet but also for deoxidation.

In the method of producing a high strength steel sheet of the present invention, when performing annealing at a low temperature of 760 ° C. to Ac ₃ ° C. at the time of heating when Tsuban a continuous annealing line or a continuous galvanizing line, ferrite and austenite during annealing The volume ratio changes depending on the annealing temperature. That is, when the annealing temperature is as low as less than 760 ° C., it takes a long time to dissolve cementite, so that austenite (martensite after cooling) becomes too small, and the maximum tensile strength of 900 MPa or more cannot be secured. On the other hand, when the annealing temperature exceeds Ac ₃ ° C, it becomes an austenite single phase region annealing, and when cooled as it is, it becomes a martensite single phase structure and the structure of the present invention cannot be obtained. For this reason, in order to manufacture a steel sheet with small material variations, it is desirable that the Ac ₃ transformation point is sufficiently high and the steel sheet structure does not change much even if the annealing temperature changes. Si and Al, since increasing the Ac ₃ transformation point, and a large amount by containing Si and Al in the above-mentioned range by increasing the Ac ₃ transformation point, be difficult to rely on steel sheet structure to the annealing temperature Is desirable.

In the present invention, Ac ₃ means the lower limit temperature at which an austenite single phase is obtained, and can be measured by investigating the temperature dependence of the length change. That is, in steel, ferrite that is stable at room temperature and austenite that is stable at high temperature have different densities and thermal expansion coefficients. As a result, the Ac ₃ transformation point can be easily measured by measuring the temperature dependence of the change in length of the test piece.

“Mn: 1.5 to 3.0%”
Mn is contained to increase the strength of the high-strength steel plate. However, if the Mn content exceeds 3.0%, the volume ratio of martensite increases too much, the volume ratio of ferrite contributing to securing ductility becomes insufficient, and the ductility and bendability become insufficient. If the Mn content exceeds 3.0%, the hardness distribution of the steel sheet surface layer due to segregation of Mn also increases. On the other hand, if the Mn content is less than 1.5%, the pearlite transformation that occurs during the cooling process cannot be suppressed, and the steel sheet structure becomes a ferrite and pearlite structure, resulting in insufficient strength. The Mn content is preferably in the range of 1.6 to 2.8%, more preferably in the range of 1.7 to 2.6%.

“Ti: 0.005 to 0.09%”
Ti is a strengthening element, and contributes to an increase in the strength of the steel sheet by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. Ti is also contained to suppress B from becoming a nitride. Although B contributes to the structure controllability at the time of hot rolling, the structure control in the continuous annealing equipment and the continuous hot dip galvanizing equipment, and high strength, this effect cannot be obtained when B becomes a nitride. However, if the Ti content exceeds 0.09%, the precipitation of carbonitrides increases and the formability deteriorates. Moreover, if the Ti content exceeds 0.09%, the recrystallization of ferrite will be significantly delayed during continuous annealing and production in continuous hot dip galvanizing equipment, so unrecrystallized ferrite remains after annealing. Easy and may result in a significant increase in yield stress. Further, Ti is preferably contained in an amount of 0.015% or more. In this case, the growth of martensite lath can be sufficiently suppressed.

“B: 0.0001 to 0.01%”
Since B delays the ferrite transformation from austenite, it can be utilized for increasing the strength of the steel sheet. In addition, since B delays the ferrite transformation from austenite during hot rolling, the hot rolled sheet can have a bainite single-phase structure, can improve the homogeneity of the hot rolled sheet, and can improve bendability. If the B content is less than 0.0001%, sufficient effects cannot be obtained. When the B content exceeds 0.01%, not only the effect is saturated, but also the productivity at the time of hot rolling is lowered. The content of B is preferably in the range of 0.0003 to 0.007%, and more preferably in the range of 0.0005 to 0.005%.

“P: 0.001 to 0.03%”
P: P tends to segregate in the central part of the plate thickness of the steel sheet and embrittles the weld. When the P content exceeds 0.03%, embrittlement of the welded portion becomes remarkable, so the appropriate range is limited to 0.03% or less. Although the lower limit of the content of P is not particularly defined, setting it to less than 0.001% is economically disadvantageous, so this value is set as the lower limit.

“S: 0.0001 to 0.01%”
S adversely affects weldability, manufacturability during casting and hot rolling. For this reason, the upper limit of the S content is set to 0.01% or less. The lower limit value of the S content is not particularly defined, but if it is less than 0.0001%, it is economically disadvantageous, so this value is the lower limit value. In addition, S combines with Mn to form coarse MnS, thereby reducing bendability. Therefore, it is necessary to reduce the S content as much as possible in order to improve bendability.

“N: 0.0005 to 0.0100%”
N forms coarse nitrides and degrades bendability and hole expansibility, so the content needs to be suppressed. This tendency becomes significant when the N content exceeds 0.0100%. In addition, if the N content exceeds 0.0100%, blowholes are generated during welding. Therefore, it is better that the N content is small. Although the lower limit of the N content is not particularly defined, the effects of the present invention are exhibited. However, if the N content is less than 0.0005%, the manufacturing cost is significantly increased. This is a reasonable lower limit.

“O: 0.0005 to 0.007%”
Since O forms an oxide and degrades bendability and hole expandability, the content needs to be suppressed. In particular, oxides often exist as inclusions, and if they are present on the punched end face or cut face, notch scratches or coarse dimples are formed on the end face, causing stress concentration during bending or strong processing. As a starting point of crack formation, the hole expandability or bendability is greatly deteriorated. This tendency becomes significant when the O content exceeds 0.007%. When the O content is less than 0.0005%, an excessive cost increase is caused, which is not economically preferable. However, even if the O content is less than 0.0005%, it is possible to ensure the tensile maximum strength of 900 MPa or more, excellent ductility and bendability, which are the effects of the present invention.

The high-strength steel sheet of the present invention may further contain the following elements as necessary.
“Nb: 0.005 to 0.09%”
Nb is a strengthening element and, like Ti, contributes to increasing the strength of the steel sheet by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. However, if the Nb content exceeds 0.09%, carbonitride precipitation increases and the formability deteriorates. In addition, if the Nb content is large, the recrystallization of ferrite is greatly delayed during production in continuous annealing and continuous hot dip galvanizing equipment, so unrecrystallized ferrite tends to remain after annealing, resulting in significant yield. This leads to an increase in stress. For this reason, the upper limit of the Nb content is preferably 0.09%. Moreover, the said effect acquired by containing Nb will become inadequate that content of Nb is less than 0.005%. The Nb content is preferably in the range of 0.01 to 0.08%, and more preferably in the range of 0.015 to 0.07%.

1 of “Cr: 0.01 to 2.0%”, “Ni: 0.01 to 2.0%”, “Cu: 0.01 to 2.0%”, “Mo: 0.01 to 0.8%” Species or two or more types Cr, Ni, Cu, and Mo are elements that contribute to the improvement of strength, and can be used in place of part of Mn. It is preferable that Cr, Ni, Cu, and Mo contain 0.01% or more of one or more of each. On the other hand, if the content of each element is too large, pickling properties, weldability, hot workability and the like may deteriorate, so the content of Cr, Ni, and Cu may be 2.0% or less. The Mo content is preferably 0.8% or less.

"V: 0.005-0.09%"
V is a strengthening element and, like Ti and Nb, contributes to increasing the strength of the steel sheet by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. Moreover, since the delayed fracture characteristic can be improved by containing V, it is desirable to contain it in the present invention. However, if the V content exceeds 0.09%, carbonitride precipitation increases and the formability deteriorates. In addition, when the V content is large, recrystallization of ferrite is greatly delayed during continuous annealing and production in continuous hot dip galvanizing equipment, so unrecrystallized ferrite tends to remain after annealing, resulting in significant yield. This leads to an increase in stress. For this reason, the upper limit of the V content is preferably 0.09%. Moreover, the said effect obtained by containing V will become inadequate that content of V is less than 0.005%. The content of V is preferably in the range of 0.01 to 0.08%, and more preferably in the range of 0.015 to 0.07%.

“Total of 0.0001 to 0.5% of one or more of Ca, Ce, Mg, and REM”
In the high-strength steel sheet of the present invention, one or more selected from Ca, Ce, Mg, and REM can be contained in a total amount of 0.0001 to 0.5%. Ca, Ce, Mg, and REM are elements used for controlling the form of oxides and sulfides. By containing one or two or more in total of 0.0001% or more, the oxide size after deoxidation In addition, it is possible to reduce the size of the sulfide present in the hot-rolled sheet, which contributes to improvement of bendability. However, if their content exceeds 0.5% in total, it will cause deterioration of molding processability.
REM is an abbreviation for Rare Earth Metal and refers to an element belonging to the lanthanoid series. In the present invention, REM and Ce are often added by misch metal and may contain a lanthanoid series element in combination with La and Ce. Even if these lanthanoid series elements other than La and Ce are included as inevitable impurities, the effect of the present invention is exhibited. Even if the metal La or Ce is added, the effect of the present invention is exhibited.

  The high-strength steel plate of the present invention may be a high-strength galvanized steel plate by forming a galvanized layer or an alloyed galvanized layer on the surface. Since the galvanized layer is formed on the surface of the high-strength steel plate, the steel sheet has excellent corrosion resistance. In addition, since the alloyed galvanized layer is formed on the surface of the high-strength steel sheet, it has excellent corrosion resistance and excellent paint adhesion.

Next, the manufacturing method of the high strength steel plate of this invention is demonstrated.
In order to manufacture the high-strength steel sheet of the present invention, first, a slab having the above-described chemical component (composition) is cast.
Next, the cast slab is directly or once cooled and then heated to 1050 ° C. or higher, and hot rolling is completed at the Ar ₃ transformation point or higher, wound in a temperature range of 400 ° C. to 670 ° C., and pickled. Then, cold rolling with a rolling reduction of 40 to 70% is performed.
In this embodiment, the slab used for hot rolling is not particularly limited. That is, the slab to be used for hot rolling may be one produced by a continuous casting slab, a thin slab caster, or the like, and is compatible with a process such as continuous casting-direct rolling (CC-DR).

The slab heating temperature for hot rolling needs to be 1050 ° C. or higher. If the slab heating temperature is less than 1050 ° C., the finish rolling temperature may be lower than the Ar ₃ point, resulting in two-phase rolling of ferrite and austenite, the hot rolled sheet structure becomes a heterogeneous mixed grain structure, Even if it passes through a rolling and annealing process, a heterogeneous structure | tissue will not be eliminated but it will be inferior to ductility and bendability. In the present invention, since a large amount of alloy elements are contained in order to ensure a maximum tensile strength of 900 MPa or more after annealing, the strength during finish rolling tends to be high. When the slab heating temperature is lower than 1050 ° C., the finish rolling temperature is lowered, the rolling load is further increased, rolling may be difficult, and the shape of the steel sheet after rolling may be deteriorated. Further, the upper limit of the slab heating temperature is not particularly defined, and the effect of the present invention is exhibited. However, it is not economically preferable to make the heating temperature excessively high. For this reason, the upper limit of the slab heating temperature is preferably less than 1300 ° C.

In the present embodiment, the Ar ₃ transformation point is calculated by the following formula.
Ar ₃ transformation point (° C.) = 901-325 × C + 33 × Si-92 × (Mn + Ni / 2 + Cr / 2 + Cu / 2 + Mo / 2)
(C, Si, Mn, Ni, Cr, Cu, and Mo in the formula are the content [% by mass] of each component in the steel.)
The hot rolling finish rolling temperature (hot rolling completion temperature) may be at least the Ar ₃ transformation point, and the effect of the present invention is exhibited without any particular upper limit. If the rolling temperature is lower than the Ar ₃ transformation point, the rolling load becomes excessively high and production becomes difficult, and hot rolling is performed in the two-phase region of ferrite and austenite. The tissue becomes uneven. That is, the ferrite formed during finish rolling is elongated by rolling and becomes coarse, and the ferrite transformed from austenite after rolling becomes fine. Such a non-uniform microstructure is not preferable because the material varies even if the microstructure is controlled by performing cold rolling-annealing. On the other hand, when the hot rolling finish rolling temperature is excessively high, the heating temperature of the slab must be excessively high in order to secure the temperature, which is not preferable. From this, it is desirable that the upper limit temperature of the finish rolling temperature of hot rolling be 1000 ° C. or less.

Next, the steel plate that has been hot-rolled is wound up in a temperature range of 400 ° C to 670 ° C. When the coiling temperature exceeds 670 ° C., coarse ferrite or pearlite structure is present in the hot rolled structure. For this reason, the structure heterogeneity after annealing becomes large, the block size exceeds 1 μm, and the bendability of the final product deteriorates. On the other hand, when the coiling temperature exceeds 670 ° C., the thickness of the oxide formed on the surface of the steel sheet is excessively increased.
Moreover, it is more preferable that the coiling temperature is 630 ° C. or less because the structure after annealing can be made fine to improve the strength ductility balance, and the structure after annealing can be uniformly dispersed to improve the bendability. However, when the coiling temperature is less than 400 ° C., the hot-rolled sheet strength is extremely increased, so that troubles such as sheet breakage and shape defects are easily induced during cold rolling. Therefore, the lower limit of the winding temperature needs to be 400 ° C.

  Note that the finish rolling may be continuously performed by joining the rough rolled plates during hot rolling. The rough rolled plate may be wound up once.

  Next, hot rolling is completed and the wound steel sheet is pickled. By pickling, the oxide on the steel sheet surface can be removed. For this reason, pickling is important in order to improve the chemical conversion property of the cold-rolled high-strength steel sheet as the final product and the hot-plating property of the cold-rolled steel sheet for hot-dip zinc or alloyed hot-dip galvanized steel sheet. Pickling may be performed once or may be performed in a plurality of times.

  Next, the hot-rolled steel sheet after pickling is cold-rolled with a rolling reduction of 40 to 70%. When the rolling reduction here is less than 40%, it becomes difficult to keep the shape of the cold-rolled steel sheet obtained after cold rolling flat, and the ductility of the final product becomes poor. On the other hand, if the rolling reduction exceeds 70%, the cold rolling load becomes too large, and cold rolling becomes difficult. The rolling reduction is more preferably 45 to 65%. In addition, the effect of the present invention is exhibited without any particular limitation on the number of rolling passes and the rolling reduction for each rolling pass.

Thereafter, the obtained cold-rolled steel sheet is passed through a continuous annealing line to produce a high-strength cold-rolled steel sheet. At this time, it is performed under the first condition or the second condition described below.
"First condition"
When passing through a continuous annealing line, the temperature between 550 ° C. and 760 ° C. is retained for 30 seconds or more during heating, and annealing is performed at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., and then an average cooling is performed between the maximum heating temperature and 630 ° C. Cool at a rate of 5 ° C./sec or less, cool between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./sec or more, and bend-bend back with a roll having a roll diameter of 350 mm or less at a temperature range of 700 ° C. to 400 ° C. After the deformation, it is cooled to room temperature at an average cooling rate of 3 ° C./second or more.

"Second condition"
When passing through the continuous annealing line, after performing annealing in the same manner as the first condition described above, the maximum heating temperature between 630 ° C and 630 ° C to 570 ° C is cooled in the same manner as the first condition, Bending-back bending deformation is performed in the same manner as in the first condition, and after holding at a temperature range of 450 ° C. to 250 ° C. for 30 seconds or more, it is cooled to room temperature at an average cooling rate of 3 ° C./second or more.
Furthermore, in this invention, it is good also as a high intensity | strength galvanized steel sheet by giving zinc electroplating to the high intensity | strength cold-rolled steel sheet obtained by letting a continuous annealing line pass in 1st condition or 2nd condition. .

In the present invention, the cold-rolled steel sheet obtained by the above method may be passed through a continuous hot dip galvanizing line to produce a high-strength galvanized steel sheet. In this case, the third condition or the fourth condition shown below is performed.
"Third condition"
When passing through the continuous hot dip galvanizing line, after annealing in the same manner as the above-mentioned first condition, cooling between the highest heating temperature and 630 ° C. is performed in the same manner as in the first condition, and then from 630 ° C. to [(Zinc plating bath temperature −40 ° C.) to (Zinc plating bath temperature + 50 ° C.)] was cooled at an average cooling rate of 3 ° C./second or more, and bending-bending deformation was performed in the same manner as in the first condition. Then, it is immersed in a galvanizing bath and cooled.
Thus, cooling between 630 ° C. to [(Zinc plating bath temperature −40 ° C.) to (Zinc plating bath temperature + 50 ° C.)] ° C. at a suitable temperature and immersing in the galvanizing bath results in zinc on the surface. A high-strength galvanized steel sheet with a plated layer is obtained.

"4th condition"
When passing through the continuous hot dip galvanizing line, after performing the steps up to immersing in the galvanizing bath in the same manner as the third condition described above, the alloying treatment is performed at a temperature of 460 ° C. to 600 ° C. and cooling is performed. I do.
Here, when the alloying treatment is performed, a Zn-Fe alloy formed by alloying a galvanized layer on the surface is formed, and a high-strength galvanized steel sheet having the alloyed galvanized layer on the surface is obtained.

In the method for producing a high-strength cold-rolled steel sheet or high-strength galvanized steel sheet according to the present invention, Nb, which contains Ti as an element for precipitating fine precipitates as a slab, and precipitates fine precipitates as necessary. , A slab having the above-described chemical components (composition) including V and the like. Therefore, when a cold-rolled steel sheet is manufactured by the above method and a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet is manufactured under any one of the first condition to the fourth condition, in the annealing process after cold-rolling Fine precipitates such as Ti are deposited. More specifically, when the cold-rolled steel sheet obtained after cold rolling is heated when passing through a continuous annealing line or a continuous hot-dip galvanizing line or during annealing near the maximum heating temperature, fine precipitation such as Ti Things are deposited. In the present invention, between 550 ° C. and 760 ° C. is retained for 30 seconds or more during the above heating, and annealing is performed at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., so fine fine precipitates such as Ti are deposited. The

In the present invention, since fine precipitates such as Ti, Nb, and V are deposited in the annealing process after cold rolling, the growth of martensite lath is suppressed as shown below, and the martensite block size is reduced. Is reduced, and a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet is obtained.
That is, since martensite is transformed from austenite, in order to suppress the growth of lath of martensite, it is preferable to disperse fine precipitates such as Ti in the austenite during annealing. In this invention, when letting a continuous annealing line or a continuous hot dip galvanizing line pass through, between 550 ° C. and 760 ° C. is retained for 30 seconds or more during heating, and annealing is performed at a maximum heating temperature of 760 ° C. to Ac ₃ ° C. Therefore, fine precipitates such as Ti are precipitated in the ferrite during heating, and ferrite containing cementite and fine precipitates such as Ti is transformed into austenite at a temperature equal to or higher than the Ac ₃ transformation point. Fine precipitates such as Ti are dispersed in the austenite during annealing.

  On the other hand, for example, when fine precipitates such as Ti, Nb, and V are deposited at the time of hot rolling and winding after completion of hot rolling, Nb, Ti, etc. The carbonitriding compound will precipitate as fine precipitates, and the fine precipitates will become coarser than when fine precipitates are deposited in a steel sheet containing many dislocations after cold rolling. In addition, when fine precipitates such as Ti, Nb, and V are deposited at the time of hot rolling and winding after completion of hot rolling, the fine precipitates such as Ti are coarsened in the annealing process after cold rolling. The effect of suppressing the growth of martensite lath is greatly reduced.

  However, when the slab contains a large amount of elements that precipitate fine precipitates such as Ti, Nb, and V, even if the fine precipitates are precipitated during the hot rolling stage or winding after completion of hot rolling, Lass growth can be sufficiently suppressed. Specifically, for example, when Ti is contained in an amount of 0.015% or more, even if fine precipitates are deposited at the time of hot rolling and winding after completion of hot rolling, the growth of martensite lath is caused. It can be sufficiently suppressed.

  In the production method of the present invention, the method for precipitating fine fine precipitates is significantly different from the conventional ferrite crystal grain refinement method described below. That is, in the conventional ferrite crystal grain refinement method, during recrystallization of ferrite, high-temperature heating is performed without causing much recovery of dislocations, so that many recrystallized ferrite nuclei are generated and ferrite is refined. It was. In such a technique, ferrite can be refined, but it is difficult to refine martensite caused by reverse transformation, and it is also difficult to control the block size constituting each martensite. Therefore, the conventional ferrite crystal grain refinement method can enjoy the advantages of higher strength due to ferrite grain refinement, but the strength of the martensite structure cannot be increased, and yield associated with ferrite grain refinement is reduced. An increase in stress is inevitable. Therefore, it has been difficult to achieve both excellent hydrogen embrittlement resistance and high strength with the conventional ferrite grain refinement method.

  Furthermore, in the manufacturing method of the high-strength cold-rolled steel sheet or high-strength galvanized steel sheet according to the present invention, the roll diameter is 350 mm or less in the temperature range of 700 ° C. to 400 ° C. in any of the first condition to the fourth condition. Since the bending-bending-back deformation is performed by the rolls of the austenite, the austenite is processed by applying strain by the bending-bending-back deformation, and dislocations are introduced as nucleation sites in the steel sheet after annealing. . Since the introduction of dislocations promotes the formation of lath having a plurality of different orientations, the martensite block size is more effectively reduced.

  Note that it is difficult to reduce the size of the block, which is an aggregate of laths, to 1 μm or less just by suppressing the growth of martensite laths forming individual blocks. Martensite lath is formed with defects (dislocations) in austenite as nuclei. The lath thus formed introduces dislocations into the surrounding austenite, and the introduced dislocations act as new martensite lath nuclei. As a result, a huge block of martensite laths having the same orientation is formed. Therefore, even if the growth of martensite laths forming individual blocks is suppressed, it is difficult to reduce the size of the blocks, which are aggregates of laths, to 1 μm or less.

  On the other hand, in the present invention, in order to reduce the block size, the growth of martensite laths forming individual blocks is suppressed, and laths having a plurality of different directions are generated. A high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance with a block size of 1 μm or less is obtained.

In addition, when manufacturing a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet using a continuous annealing line or a continuous hot-dip galvanizing line, for example, even if a large amount of dislocations are introduced during cold rolling, an annealing process after cold rolling Recrystallization or reverse transformation causes dislocations to disappear and new austenite is formed. Therefore, it is not preferable to introduce nucleation sites for forming laths having a plurality of different orientations before annealing.
On the other hand, in the manufacturing method of the high-strength cold-rolled steel sheet or the high-strength galvanized steel sheet according to the present invention, after annealing by bending-bending-back deformation in any one of the first condition to the fourth condition. Since nucleation sites are introduced into the steel sheet, the formation of laths having a plurality of different orientations is effectively promoted by the nucleation sites.

In the present invention, a nucleation site is introduced into the steel sheet after annealing by bending-bending deformation, so that it is difficult to reduce the thickness and is preferable.
On the other hand, for example, when strain is introduced by tensile tension, the plate thickness and width become non-uniform due to the reduction in plate thickness, and in extreme cases, there is a possibility that the plate breaks in the furnace due to necking. Therefore, it is desirable to perform bending-bending unbending deformation.
In addition, when performing bending-bending unbending deformation, it is possible to easily introduce plastic deformation by applying a tensile stress to the plate. Therefore, a tensile stress may be applied to the plate when bending-unbending deformation is performed as long as it does not impair the plate passing property.

  In the present invention, since bending-bending deformation is performed in a temperature range of 700 ° C. to 400 ° C., recovery of dislocation as a nucleation site is suppressed, and a nucleation site can be reliably introduced. . When the temperature exceeds 700 ° C., even if strain is introduced into the austenite by bending-bending deformation, the introduced dislocation may recover and the dislocation may not act as a nucleation site. In addition, when the temperature is less than 400 ° C., the yield strength in the hot state of the steel sheet is too high, and sufficient plastic deformation cannot be performed, and the introduction of dislocations as nucleation sites that leads to refinement of the block size. May not be possible.

In the present invention, it is necessary to perform bending-bending deformation with a roll having a roll diameter of 350 mm or less. When the roll diameter exceeds 350 mm, sufficient plastic deformation cannot be performed, and dislocations cannot be introduced into austenite necessary for miniaturization of the block size. For this reason, the roll diameter needs to be 350 mm or less. Moreover, in order to add a larger amount of deformation, the roll diameter is preferably as small as possible.
The pushing amount of the roll is 1 mm or more, desirably 2 mm or more. As the amount of indentation of the roll increases, the material undergoes large deformation, and dislocations that become martensite nucleation sites are introduced into the austenite. However, there is a concern that an extreme increase in the amount of pressing may cause deterioration of sheet passing properties and formation of wrinkles on the steel sheet surface. For this reason, it is desirable that the indentation amount be a minimum value that allows the introduction of strain into the austenite. For example, it is preferable that the indentation amount be 700 mm or less.

In the present invention, after annealing, between the maximum heating temperature and 630 ° C. is cooled at an average cooling rate of 14 ° C./second or less, the formation of ferrite is promoted, and the volume fraction of ferrite contributing to securing ductility is set. The yield stress (YS) is small and the yield ratio (YR) is 0.75 or less. Therefore, the residual stress is small, and it has excellent delayed fracture resistance. In the present invention, since the ferrite volume fraction can be 50% or more, C is concentrated in martensite, and the C concentration in martensite is 0.3% to 0.9%. , High strength martensite is obtained.
On the other hand, when cooling between the maximum heating temperature and 630 ° C. at an average cooling rate exceeding 5 ° C./sec, the ferrite volume fraction becomes less than 50%, or the yield ratio exceeds (YR) 0.75. As a result, at least one of the ductility, delayed fracture resistance, and strength of the high-strength cold-rolled steel sheet or high-strength galvanized steel sheet may be insufficient.

In the present invention, the average cooling rate between 630 ° C. and 570 ° C. or between 630 ° C. and [(Zinc plating bath temperature−40 ° C.) − (Zinc plating bath temperature + 50 ° C.)] ° C. is 3 ° C./second or more. Therefore, pearlite and bainite transformation that can occur in this temperature range can be suppressed, and austenite can be efficiently transformed into martensite.
On the other hand, when the average cooling rate in the above temperature range is less than 3 ° C./second, the austenite is transformed into pearlite, the martensite volume fraction is insufficient, and the high-strength cold-rolled steel sheet or high-strength zinc The strength of the plated steel sheet may be insufficient. In addition, it is not preferable that the average cooling rate in the above temperature range is less than 3 ° C./second because productivity is lowered.

Moreover, in this invention, the atmosphere in the annealing furnace of the continuous annealing line or the continuous hot dip galvanizing line at the time of manufacture of a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet contains ₁ to 60% by volume of H2. The atmosphere is composed of the balance N ₂ , H ₂ O, O ₂ and inevitable impurities, and the logarithm log (PH ₂ O / PH ₂ ) of the water pressure and the hydrogen partial pressure in the atmosphere is −3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5 is preferable.

  When the atmosphere in the annealing furnace is the atmosphere described above, before Si, Mn, and Al contained in the steel sheet diffuse into the steel sheet surface, O diffused into the steel sheet and Si, Mn, and Al contained in the steel sheet are contained. As a result of the reaction, oxides are formed inside the steel sheet, and the formation of oxides composed of these elements on the steel sheet surface is suppressed. Therefore, by setting the atmosphere in the annealing furnace to the atmosphere described above, non-plating due to the formation of oxide on the steel sheet surface can be suppressed and the alloying reaction can be promoted. It is possible to prevent deterioration of chemical conversion properties due to the formation of oxides.

  In addition, the ratio of the water pressure and the hydrogen partial pressure in the atmosphere in the annealing furnace can be adjusted by a method of blowing water vapor into the furnace. Thus, the method of adjusting the ratio between the water pressure and the hydrogen partial pressure in the atmosphere in the annealing furnace is simple and preferable.

Further, in an atmosphere of the annealing furnace, the concentration of H ₂ exceeds 60 vol% is not preferable because it increases the cost. On the other hand, when the H ₂ concentration is less than 1% by volume, Fe contained in the steel plate is oxidized, and there is a possibility that wettability and plating adhesion are insufficient.
In addition, by setting the logarithm log (PH ₂ O / PH ₂ ) of moisture pressure and hydrogen partial pressure in the atmosphere in the furnace to −3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5, Even if it is a steel containing a large amount, sufficient plating properties can be secured. Note that the lower limit of the logarithm log (PH ₂ O / PH ₂ ) of the water pressure and the hydrogen partial pressure was set to −3 or more. This is because there is a risk that the wettability and plating adhesion may be reduced. On the other hand, the upper limit of the logarithm log (PH ₂ O / PH ₂ ) of the water pressure and the hydrogen partial pressure is set to −0.5 because the effect is saturated.

On the other hand, when manufacturing the high-strength cold-rolled steel sheet or high-strength galvanized steel sheet according to the present invention using the conventional steel sheet manufacturing method, the atmosphere in the annealing furnace at the time of manufacturing is set as the atmosphere in the annealing furnace. As a result, the following problems may occur.
That is, in the present invention, the above-described chemical component (composition) including Si (or Si and Al) to improve the ferrite volume fraction and ensure ductility as the slab, and Mn to increase the strength of the high-strength steel sheet. The slab which has is used. Since Si, Mn, and Al are elements that are very easily oxidized compared to Fe, even in a reducing atmosphere of Fe, the surface of the steel sheet containing Si (or Si and Al) and Mn has Si Oxides (or Si oxides and Al oxides) and Mn oxides are formed.

Oxides containing Si, Mn, and Al alone or in combination formed on the surface of the steel sheet cause deterioration of the chemical conversion property in the high-strength cold-rolled steel sheet. In addition, these oxides have poor wettability with molten metals such as zinc, and therefore cause non-plating when forming a galvanized layer on the surface of high-strength steel sheets containing Si (or Si and Al). It becomes. Si and Al may cause problems such as delaying alloying when producing a high-strength galvanized steel sheet subjected to alloying treatment.
Here, as a method for suppressing the formation of oxides on the surface of the steel sheet, a method in which the atmosphere in the annealing furnace is made a reducing atmosphere of each element is considered. In this embodiment, the atmosphere in the annealing furnace is changed to the atmosphere in the annealing furnace. Although the above-described atmosphere was used, an atmosphere such as Si, Mn and Al was easily oxidized although it was a reducing atmosphere of Fe.

The present invention is not limited to the above example.
For example, in the above-described method for producing a high-strength cold-rolled steel sheet or a high-strength galvanized steel sheet, the moisture pressure and hydrogen partial pressure are controlled to control the atmosphere in the annealing furnace. You may control the atmosphere in an annealing furnace using the method of controlling a pressure, the method of controlling the partial pressure of nitrogen dioxide and nitric oxide, or the method of blowing oxygen directly in a furnace. Even in this case, Si, Mn, and Al are contained alone or in combination in the steel plate near the surface layer, as in the case where the atmosphere in the annealing furnace is controlled by controlling the moisture pressure and the hydrogen partial pressure. An oxide can be deposited, and the same effect as described above can be obtained.

  In the method for producing a high-strength galvanized steel sheet according to the present invention, in order to improve plating adhesion, the steel sheet before annealing is plated with one or more kinds selected from Ni, Cu, Co and Fe. May be.

Moreover, when manufacturing the high-strength galvanized steel sheet of the present invention, as a process from annealing to dipping in a galvanizing bath, “after degreasing and pickling, heating in a non-oxidizing atmosphere, including H ₂ and N ₂ . After annealing in a reducing atmosphere, cool it down to near the galvanizing bath temperature and immerse it in the galvanizing bath ”or“ Adjust the atmosphere during annealing, first oxidize the steel plate surface, then reduce After cleaning the surface of the steel sheet before plating, it is immersed in a galvanizing bath, or the total reduction furnace method, or "after the steel sheet is degreased and pickled, and flux treatment is performed using ammonium chloride, A flux method or the like that is then immersed in a galvanizing bath may be used.

"Example 1"
The cast slab having the chemical composition (composition) shown in Table 1 is directly heated to the temperature (slab heating temperature) shown in Table 2 to Table 5, and the temperature shown in Table 2 to Table 5 (hot rolling completion temperature). The hot rolling was completed at the temperature range shown in Tables 2 to 5 and wound up in the temperature range (winding temperature). After pickling, cold rolling was performed at the rolling reductions shown in Tables 2 to 5, and Experimental Examples 1 to The steel plate of Experimental Example 99 was obtained.

  In Tables 2 to 5, CR, which indicates the type of product plate, is a cold-rolled steel sheet manufactured by a continuous annealing line, EG is an electroplated steel sheet obtained by electroplating zinc on the cold-rolled steel sheet CR, and GI is a continuous melt. A galvanized steel sheet manufactured by a galvanizing line, GA is an alloyed galvanized steel sheet manufactured by a continuous hot dip galvanizing line.

Subsequently, some of the experimental examples among the steel plates of Experimental Examples 1 to 99 were passed through a continuous annealing line to produce a cold-rolled steel plate CR. The atmosphere in the annealing furnace of the continuous annealing line is an N ₂ atmosphere containing 1% by volume of H _2, and the logarithm log (PH ₂ O / PH ₂ ) of the moisture pressure and the hydrogen partial pressure in the atmosphere in the furnace is − 2.8.

  In addition, when passing through the continuous annealing line, during the heating, between 550 ° C. and 760 ° C. was retained at the residence time shown in Tables 6 to 9, and after annealing at the maximum heating temperature shown in Tables 6 to 9, The temperature between the maximum heating temperature and 630 ° C. is cooled at the average cooling rate shown in Tables 6 to 9, the temperature between 630 ° C. and 570 ° C. is cooled at the average cooling rate shown in Tables 6 to 9, and the temperature is 700 ° C. to 400 ° C. In the range, the rolls having the roll diameters shown in Table 6 to Table 9 were subjected to bending-bending deformation with the pushing amount shown in Tables 6 to 9. Thereafter, in some experimental examples, the holding temperatures shown in Tables 6 to 9 were held at the holding times shown in Tables 6 to 9, and then cooled to room temperature at the average cooling rates shown in Tables 6 to 9. In other experimental examples, it was cooled to room temperature.

Then, about the one part of the experimental example which let the continuous annealing line pass, the zinc system electroplating was performed with the method shown below, and the electroplating steel plate EG was obtained.
First, as a pretreatment for plating, alkaline degreasing, water washing, pickling, and water washing were sequentially performed on a steel plate that was passed through a continuous continuous annealing line. Thereafter, a liquid circulation type electroplating apparatus is used for the steel sheet after the pretreatment, and a plating bath made of zinc sulfate, sodium sulfate, and sulfuric acid is used until a predetermined plating thickness is obtained at a current density of 100 A / dm ^2. Electrolytic treatment was performed and Zn plating was performed.

Moreover, the cold-rolled steel plate which was not made to let the continuous annealing line pass among the steel plates of Experimental example 1-Experimental example 99 was made to pass a continuous hot dip galvanizing line, and the galvanized steel plate GI was manufactured. The atmosphere in the annealing furnace of the continuous hot dip galvanizing line is an N ₂ atmosphere containing 1% by volume of H _2, and the logarithm log of the water pressure and the hydrogen partial pressure in the furnace atmosphere (PH ₂ O / PH ₂ ). Was -1.2.

When passing through a continuous hot dip galvanizing line, during the heating, between 550 ° C. and 760 ° C. was retained for the residence time shown in Table 6 to Table 9, and after annealing at the maximum heating temperature shown in Table 6 to Table 9, Cooling between the highest heating temperature and 630 ° C. is performed at an average cooling rate shown in Tables 6 to 9, and cooling between 630 ° C. and galvanizing bath temperature is performed at an average cooling rate shown in Tables 6 to 9, and 700 ° C. to 400 ° C. After performing the bending-bending deformation of the indentation amounts shown in Tables 6 to 9 with rolls having the roll diameters shown in Tables 6 to 9 in the temperature range, the samples were immersed in a galvanizing bath and cooled.
Moreover, in some experimental examples, after performing the process until it is immersed in a galvanizing bath, alloying treatment is performed at the temperatures shown in Tables 6 to 9, and the average cooling rates shown in Tables 6 to 9 are used. Cooling was performed.

  For the high-strength cold-rolled steel sheets or high-strength galvanized steel sheets of Experimental Examples 1 to 99 thus obtained, the steel sheet structure was observed by the FE-SEM-EBSP method, and the volume of ferrite, martensite, and remaining structure And the average crystal grain size of ferrite and martensite were examined. The results are shown in Tables 10 to 13.

  In Tables 10 to 13, P described in the remaining structure means pearlite, B means bainite, RA means retained austenite, and C means cementite.

Moreover, the block size of the martensite was measured by the method shown below about the high-strength cold-rolled steel plate or the high-strength galvanized steel plate of Experimental Example 1 to Experimental Example 99. The results are shown in Tables 10 to 13.
For the measurement of the block size of martensite, the FE-SEM-EBSP method is used, and by conducting a preliminary experiment, the block size of several martensite grains is measured and the orientation of the structure can be analyzed. Measurements were taken after the interval was determined to be 25 nm.

Moreover, about the high intensity | strength cold-rolled steel plate or high-strength galvanized steel plate of Experimental example 1-Experimental example 99, the C density | concentration in a martensite was calculated | required by the method shown below.
First, a steel plate was embedded and polished in a cross section parallel to the rolling direction. Then, the structure | tissue observation by SEM was performed and the position of each structure | tissue was pinpointed. Then, using EPMA, the area of 50 μm length and 50 μm width of the steel sheet was subjected to surface analysis with 10 visual fields and step size of 0.1 μm, and the average value of C concentration contained in each martensite grain obtained was It was defined as the C concentration in martensite.
The results are shown in Tables 10 to 13.

Further, a tensile test piece based on JIS Z 2201 was taken from the high-strength cold-rolled steel sheet or the high-strength galvanized steel sheet of Experimental Examples 1 to 99, and the tensile test was performed based on JIS Z 2241. (TS), yield stress (YS), elongation (El), uniform elongation (u-El), and total elongation (El) were measured.
Further, the yield ratio (YR = (YS / TS)) is calculated using the measured tensile maximum strength (TS) and yield stress (YS), and the measured tensile maximum strength (TS) and total elongation (El.). Was used to calculate the strength-ductility balance (TS × El.).

Moreover, the delayed fracture test was done by the method shown below about the high-strength cold-rolled steel plate or the high-strength galvanized steel plate of Experimental Examples 1 to 99.
First, a rectangular test piece having a length of 150 mm and a width of 100 mm cut in a direction perpendicular to the rolling direction was cut, and the end portion was mechanically ground. Thereafter, bending was performed at the center of the length in the horizontal direction (rolling direction), and the cross section in the vertical direction (perpendicular to the rolling direction) was made substantially L-shaped. Then, the test piece of the shape shown in FIG. 2 was obtained by carrying out the flange-up process of the vertical direction edge part of a test piece, making the cross section of a horizontal direction into substantially U shape.

  FIG. 2 is a view for explaining the shape of the test piece after being flange-formed, and FIG. 2 (a) is a perspective view of the test piece. FIG. 2B is a cross-sectional view showing a part of the test piece, and is a cross-sectional view showing a part subjected to bending. Moreover, FIG.2 (c) is sectional drawing which showed a part of test piece, and is sectional drawing which showed the part which performed the flange-up process.

As shown in FIG. 2 (a), the test piece after the flange molding is made up of a pair of flat plate portions 1a and 1b extending in a direction orthogonal to each other, and a curved surface connecting the flat plate portion 1a and the flat plate portion 1b. It has a substantially L-shaped main portion 1 having a bent portion 1c. As shown in FIGS. 2A and 2B, the bending portion 1c is formed by bending at a bending radius (r ₀ ) of 10 mm and a bending angle (θ _A ) of 90 °. is there.
Further, as shown in FIG. 2A, flanges 2 are formed at both edge portions in the extending direction of a line where the flat plate portion 1a and the flat plate portion 1b of the main portion 1 are orthogonal to each other. As shown in FIGS. 2 (a) and 2 (c), the flange 2 has a height (h) and a bending angle (θ _B ) of 90 ° from the main portion 1 to the opposite side to the center side in bending. It is formed by flange-up processing.

  In the flange-up process of the test piece, the following formula (1) and formula (2) are used, and the nominal strain (ε) acting on the plate end of the test piece by the flange-up process is 0.04. The height (h) of the flange was calculated (0.4 mm), and the flange was molded.

ε = (2hθ _A sin θ _B ) / πr ₀ (1)
In Equation (1), h is the height of the flange, θ _A is the bending angle of the bending process, θ _B is the bending angle of the flange-up process, and r ₀ is the bending radius of the bending process. . Here, when θ _A and θ _B are π / 2, the above formula (1) is written as the following formula (2).
ε = h / r ₀ (2)

Next, the test piece after the flange molding shown in FIG. 2 is immersed in 0.5 mol / l sulfuric acid, the steel plate side is the cathode, the platinum electrode is the anode, and the current is applied at a current density of 0.1 mA / cm ² . Then, electrolysis was performed for 2 hours by electric current, and hydrogen generation and hydrogen intrusion into the steel sheet were performed. Then, the test piece was taken out and the presence or absence of the crack in the edge part of a test piece was investigated by observing the stretch flange molding part visually. Since there is a large residual stress at the end of the test piece, its progress is rapid when a crack occurs. From this, in this example, when there was a crack, all were large opening cracks, and it was possible to easily determine the presence or absence of a crack even visually. In this example, the end of the test piece was carefully observed using a magnifying glass, an actual microscope, etc., and the presence or absence of cracks was confirmed again, and it was confirmed that those without open cracks had no fine cracks.
In the delayed fracture test results shown in Tables 10 to 13, ◯ indicates that no cracks occurred at the end portions, and x indicates that cracks occurred at the end portions. The results are shown in Tables 10 to 13.

As shown in Tables 10 to 13, in the experimental example which is an example of the present invention, the volume ratio of ferrite in the steel sheet structure was 50% or more, and it was confirmed that martensite was included.
From Table 10 to Table 13, in the experimental example which is an example of the present invention, the block size was 1 μm or less, and the C concentration in martensite was in the range of 0.3% to 0.9%.
Further, from Tables 10 to 13, in the experimental examples which are examples of the present invention, the maximum tensile strength (TS) is 900 MPa or more, the yield ratio (YR) is 0.75 or less, and the maximum tensile strength and elongation are good. Met.
Moreover, from Table 10 to Table 13, in the experimental example which is an example of the present invention, the delayed fracture test results were all “good”.

  On the other hand, in Experimental Example 2 which is a comparative example of the present invention, cooling is performed between the maximum heating temperature and 630 ° C. at an average cooling rate exceeding 5 ° C./sec. The concentration of C in martensite was insufficient, the C concentration in martensite was less than 0.3%, the yield ratio (YR) exceeded 0.75, and the elongation was insufficient. In Experimental Example 2, the delayed fracture test result was also x.

In Experimental Example 8, which is a comparative example of the present invention, since the average cooling rate between 630 ° C. and the galvanizing bath temperature is less than 3 ° C./second, the austenite is transformed into pearlite, so that the volume ratio of martensite. , The maximum tensile strength (TS) was less than 900 MPa, and the strength was insufficient.
In Experimental Example 54, which is a comparative example of the present invention, since the average cooling rate between 630 ° C. and the galvanizing bath temperature is less than 3 ° C./second, the austenite transforms into pearlite, so that the martensite volume fraction 0%, the maximum tensile strength (TS) was less than 900 MPa, the yield ratio (YR) exceeded 0.75, and the strength and elongation were insufficient.

Further, in Experimental Example 3, Experimental Example 14, Experimental Example 19, Experimental Example 40, and Experimental Example 50, which are comparative examples of the present invention, the maximum heating temperature during annealing is less than 760 ° C., so the volume of martensite The rate was 0%, the maximum tensile strength (TS) was less than 900 MPa, and the strength was insufficient.
Moreover, in Experimental example 24, Experimental example 29, Experimental example 33, Experimental example 46, Experimental example 66, Experimental example 75, and Experimental example 80 which are comparative examples of this invention, the maximum heating temperature at the time of annealing heating is less than 760 degreeC. Therefore, the volume ratio of martensite is 0%.
In Experimental Example 62, which is a comparative example of the present invention, the maximum heating temperature during annealing is less than 760 ° C., so the volume ratio of martensite is 0% and the yield ratio (YR) is 0. More than .75, the elongation was insufficient. In Experimental Example 62, the delayed fracture test result was x.

In Experimental Example 6, Experimental Example 10, and Experimental Example 15, which are comparative examples of the present invention, since the push amount of the bending-bending deformation is 0 mm, the block size exceeds 1 μm, and the maximum tensile strength (TS) is The strength was less than 900 MPa, and the strength was insufficient.
In Experimental Example 20, which is a comparative example of the present invention, since the push-in amount of bending-unbending deformation is 0 mm, the block size exceeds 1 μm, the maximum tensile strength (TS) is less than 900 MPa, and the strength is insufficient. Met.
Experimental Example 25, Experimental Example 30, Experimental Example 34, Experimental Example 36, Experimental Example 41, Experimental Example 45, Experimental Example 51, Experimental Example 61, Experimental Example 67, Experimental Example 76, Experimental Example 81, which are comparative examples of the present invention Then, since the pushing amount of bending-bending unbending deformation was 0 mm, the block size exceeded 1 μm.

  In Experimental Example 26, Experimental Example 35, Experimental Example 42, Experimental Example 49, Experimental Example 65, and Experimental Example 77, which are comparative examples of the present invention, the coiling temperature after completing the hot rolling exceeds 670 ° C. Therefore, the block size exceeded 1 μm.

In Experimental Example 83 and Experimental Example 84, which are comparative examples of the present invention, since the C concentration of steel is less than 0.07%, the C concentration in martensite is less than 0.3%, and the maximum tensile strength ( TS) was less than 900 MPa, and the strength was insufficient.
In Experimental Example 85 and Experimental Example 86, which are comparative examples of the present invention, since Ti and B are not contained in the steel, the C concentration in martensite exceeds 0.9%, and the maximum tensile strength (TS) Was less than 900 MPa, and the strength was insufficient.

In Experimental Example 87 and Experimental Example 88, which are comparative examples of the present invention, since the Si concentration of the steel was less than 0.3%, the block size exceeded 1 μm.
In Experimental Example 89, which is a comparative example of the present invention, since the Mn concentration of the steel is less than 1.5%, the pearlite transformation cannot be suppressed, the martensite volume ratio becomes 0%, and the maximum tensile strength ( TS) was less than 900 MPa, and the strength was insufficient.
In Experimental Example 90, which is a comparative example of the present invention, since the Mn concentration of the steel is less than 1.5%, the pearlite transformation cannot be suppressed, the martensite volume ratio is insufficient, and the block size is 1 μm. The tensile maximum strength (TS) was less than 900 MPa, and the strength was insufficient.

In Experimental Example 91 and Experimental Example 92, which are comparative examples of the present invention, since the Mn concentration of the steel exceeds 3.0%, the volume ratio of martensite is too large, and the volume of ferrite contributing to ensuring ductility. The ratio was insufficient, the C concentration in martensite was less than 0.3%, the block size exceeded 1 μm, the yield ratio (YR) exceeded 0.75, and the elongation was insufficient. In Experimental Example 91 and Experimental Example 92, the delayed fracture test result was also x.
In Experimental Example 93 and Experimental Example 94, which are comparative examples of the present invention, since the steel does not contain Ti, the maximum tensile strength (TS) was less than 900 MPa, and the strength was insufficient.

In Experimental Example 95 and Experimental Example 96, which are comparative examples of the present invention, since steel does not contain B, the maximum tensile strength (TS) was less than 900 MPa, and the strength was insufficient.
In Experimental Example 97 and Experimental Example 98, which are comparative examples of the present invention, since the C concentration of steel exceeds 0.25%, the volume ratio of ferrite is less than 50%, and the block size exceeds 1 μm. The yield ratio (YR) exceeded 0.75 and the elongation was insufficient. In Experimental Example 97 and Experimental Example 98, the delayed fracture test result was also x.
In Experimental Example 99, which is a comparative example of the present invention, the slag heating temperature was too low, so the block size exceeded 1 μm, the maximum tensile strength (TS) was less than 900 MPa, and the strength was insufficient.

"Example 2"
Further, with respect to the high strength cold-rolled steel sheets of Experimental Example 2, Experimental Example 23, and Experimental Example 91, the yield stress (YS) -strain curve (ε) and the maximum tensile strength (TS) are shown as follows in accordance with JIS Z 2241. Asked. That is, up to 0.8% strain after yielding, the stress increase rate is constant at 20 N / (mm ² · S), and thereafter 20 mm / min. The stress-strain curve was measured at a constant crosshead speed. However, a clear upper yield point did not appear in the stress-strain curve of Experimental Example 23, which is an example of the present invention. From this, 0.2% proof stress was measured by the offset method, and this value was determined as the yield stress.
The result is shown in FIG.

  As shown in Tables 10 and 13, the tensile strengths (TS) of Experimental Example 2, Experimental Example 23, and Experimental Example 91 were all equal to or greater than 1100 MPa, but as shown in FIG. In Experimental Example 23, which is an example, compared to Experimental Example 2 and Experimental Example 91, which are comparative examples of the present invention, the yield stress (YS) when the strain is ε1 shown in FIG. 1 is low (YS1> YS3). YS2> YS3), and the yield strength after molding is also low (σ1> σ3, σ2> σ3). Thus, in Experimental Example 23, which is an example of the present invention, the residual stress is smaller than in Experimental Example 2 and Experimental Example 91, which are comparative examples of the present invention, and is excellent in delayed fracture resistance. I understand that.

DESCRIPTION OF SYMBOLS 1 ... Main part, 1a, 1b ... Flat plate part, 1c ... Processing part, 2 ... Flange.

Claims (13)

% By mass
C: 0.07 to 0.25%,
Si: 0.3-2.50%,
Mn: 1.5-3.0%
Ti: 0.005 to 0.09%,
B: 0.0001 to 0.01%
P: 0.001 to 0.03%,
S: 0.0001 to 0.01%
Al: 2.5% or less,
N: 0.0005 to 0.0100%,
O: 0.0005 to 0.007%,
And the balance is steel consisting of iron and inevitable impurities,
The steel sheet structure is mainly composed of ferrite, contains martensite, the ferrite volume fraction is 50% or more, the martensite block size is 1 μm or less, and the C concentration in the martensite is 0.3% to 0.9%. %, Yield ratio (YR) consisting of ratio of maximum tensile strength (TS) and yield stress (YS) is 0.75 or less, good tensile maximum strength with good ductility and delayed fracture resistance A high-strength steel sheet having 900 MPa or more. Furthermore, in mass%,
Nb: 0.005 to 0.09%
The high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to claim 1. Furthermore, in mass%,
Cr: 0.01 to 2.0%,
Ni: 0.01 to 2.0%,
Cu: 0.01 to 2.0%,
Mo: 0.01 to 0.8%
The high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to claim 1 or 2, characterized by containing at least one of the following. Furthermore, in mass%,
V: 0.005-0.09%
A high-strength steel sheet having a maximum tensile strength of 900 MPa or more with good ductility and delayed fracture resistance according to any one of claims 1 to 3.   Furthermore, 0.0001-0.5% of 1 type or 2 types in total of Ca, Ce, Mg, and REM is contained by mass%, The any one of Claims 1-4 characterized by the above-mentioned. A high-strength steel sheet having a maximum tensile strength of 900 MPa or more with good ductility and delayed fracture resistance.   The high-strength steel sheet having a tensile maximum strength of 900 MPa or more with good ductility and delayed fracture resistance according to any one of claims 1 to 5, wherein the surface has a galvanized layer. It is a manufacturing method of the high intensity steel plate according to any one of claims 1 to 5 ,
A slab having the chemical component according to any one of claims 1 to 5 is cast, directly or once cooled and then heated to 1050 ° C or higher, and the hot rolling is completed at an Ar ₃ transformation point or higher, 400 ° C. When winding in a temperature range of ˜670 ° C., pickling, cold rolling with a rolling reduction of 40% to 70%, and passing through a continuous annealing line,
After heating between 550 ° C. and 760 ° C. for 30 seconds or longer and annealing at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., cooling between a maximum heating temperature of 630 ° C. and an average cooling rate of 14 ° C./sec or less is performed. , The temperature between 630 ° C. and 570 ° C. is cooled at an average cooling rate of 3 ° C./second or more, and bending-bending deformation with an indentation amount of 1 mm or more is performed using a roll having a roll diameter of 350 mm or less in a temperature range of 700 ° C. to 400 ° C. A method for producing a high-strength cold-rolled steel sheet having a maximum tensile strength of 900 MPa or more with good ductility and delayed fracture resistance.   The ductility and delay resistance according to claim 7, wherein after cooling between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more, the temperature is maintained at 450 ° C. to 250 ° C. for 30 seconds or more. A method for producing a high-strength cold-rolled steel sheet having a tensile maximum strength of 900 MPa or more with good fracture characteristics. In the annealing furnace of the continuous annealing line, an atmosphere containing ₁ to 60% by volume of H ₂ and the balance N ₂ , H ₂ O, O ₂ and unavoidable impurities is used, and the moisture pressure and hydrogen content in the atmosphere The logarithm log (PH ₂ O / PH ₂ ) of pressure is set to -3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5, and ductility and delayed fracture resistance according to claim 7 or 8 A method for producing a high-strength cold-rolled steel sheet having a maximum tensile strength of 900 MPa or more with good characteristics.   Good ductility and delayed fracture resistance characterized by applying zinc electroplating after producing a high strength cold rolled steel sheet by the method for producing a high strength cold rolled steel sheet according to any one of claims 7 to 9. A method for producing a high-strength galvanized steel sheet having a maximum tensile strength of 900 MPa or more. It is a manufacturing method of the high strength steel plate according to claim 6,
A slab having the chemical component according to any one of claims 1 to 5 is cast, directly or once cooled and then heated to 1050 ° C or higher, and the hot rolling is completed at an Ar ₃ transformation point or higher, 400 ° C. Winding in a temperature range of ˜670 ° C., pickling, cold rolling with a rolling reduction of 40 to 70%, and passing a continuous hot dip galvanizing line,
After heating at 550 ° C. to 760 ° C. for 30 seconds or more during heating and annealing at a maximum heating temperature of 760 ° C. to Ac ₃ ° C., cooling between the maximum heating temperature to 630 ° C. at an average cooling rate of 14 ° C./second or less is performed. ℃ ~ [(Zinc plating bath temperature -40 ℃) ~ (Zinc plating bath temperature + 50 ℃)] is cooled at an average cooling rate of 3 ℃ / second or more, roll diameter 350mm in the temperature range of 700 ℃ ~ 400 ℃ Using the following rolls, a bending-unbending deformation with an indentation amount of 1 mm or more is performed, followed by immersion in a galvanizing bath and cooling. A tensile strength of 900 MPa or more with good ductility and delayed fracture resistance A method for producing a high-strength galvanized steel sheet having   The maximum tensile strength of 900 MPa having good ductility and delayed fracture resistance according to claim 11, wherein the alloy is subjected to an alloying treatment at a temperature of 460 ° C. to 600 ° C. after being immersed in the zinc plating bath. The manufacturing method of the high intensity | strength galvanized steel plate which has the above. In the annealing furnace of the continuous hot dip galvanizing line, an atmosphere containing ₁ to 60% by volume of H ₂ and the balance consisting of the balance N ₂ , H ₂ O, O ₂ and unavoidable impurities, and the moisture pressure in the atmosphere, The logarithm log (PH ₂ O / PH ₂ ) of the hydrogen partial pressure is set to -3 ≦ log (PH ₂ O / PH ₂ ) ≦ −0.5. A method for producing a high-strength galvanized steel sheet having a maximum tensile strength of 900 MPa or more with good delayed fracture characteristics.

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