JP4811288B2 - High-strength cold-rolled steel sheet and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet and a method for producing the same. More specifically, the present invention is suitable as a material for structural members such as automobiles and various industrial machines, and is excellent in formability and material stability at the material stage before processing, and at the time of collision at the member stage after processing. The present invention relates to a high-strength cold-rolled steel sheet having excellent impact characteristics indicated by the energy absorption performance of the same and a method for producing the same.

CO ₂ emissions to protect the global environment For reduction, for example, automobiles need to further promote weight reduction. Moreover, from the viewpoint of improving collision safety, a structural member for an automobile needs a material having a thin plate thickness and high deformation strength. Against this background, there is a demand for expanding the application of high-strength steel sheets and increasing the strength of applied materials. Recently, there is a strong demand for the application of so-called high-strength steel sheets that exceed the tensile strength of 780 MPa and further exceed 980 MPa even in the frame that controls the crash safety of automobiles. Here, the automobile frame is a typical shock absorbing member for automobiles, and is a so-called side member. In addition, a pillar etc. are illustrated as an impact-absorbing member for motor vehicles.

  In order to satisfy these demands, steel sheets must have formability when forming parts, material stability that leads to increased precision of formed parts, especially reduced tensile strength variation, and welding during assembly. In addition, after assembly, there is a need for stable collision energy performance that absorbs collision energy by plastic deformation without causing, for example, weld fracture or base metal fracture during actual collision.

(1) As a high-strength thin steel sheet having excellent formability, a composite structure steel sheet having ferrite as a main phase and a low-temperature transformation phase such as martensite or bainite as a second phase has been proposed. Patent Document 1 discloses a hot-dip galvanized steel sheet having a composite structure having ferrite as a main phase, having a tensile strength of 780 MPa (80 kgf / mm ² ) or more and a yield ratio of 60% or less. This steel sheet has a formability of a tensile strength-elongation balance (TS × El) of about 13,000 to 25000 MPa ·%. However, in this way, the high-strength thin steel sheet using a hard low-temperature transformation phase is liable to form a crack at the interface between the hard phase and the soft phase, so that, for example, bending performance largely governed by local ductility deteriorates. . Furthermore, since the so-called HAZ softening occurs in which the hard phase of the heat-affected zone is softened during welding, there is a problem that the weld zone fracture resistance is not sufficient.

  In order to suppress the occurrence of cracks, it is necessary to have a uniform structure with a small hardness difference, and in order to make it difficult to cause HAZ softening during welding, it is necessary to suppress the use of the hard phase as much as possible. Therefore, a steel sheet has been proposed that actively utilizes precipitation strengthening that is not transformation strengthening utilizing a hard phase.

(2) Patent Document 2 discloses a high-strength, high-yield-ratio hot-dip galvanized steel sheet having a non-composite structure with a tensile strength of 440 MPa (45 kg / mm ² ) or more and a yield ratio of 80% or more. This steel sheet has a high strength with a tensile strength of 700 MPa or more and a yield ratio of 80% or more by adding Ti and Nb, which are carbonitride-forming elements, and forming a two-phase structure of ferrite and austenite phases during continuous annealing. Is shown. However, when steel added with Ti and Nb is annealed at a temperature that has a two-phase structure, it becomes a band structure, resulting in a large variation in mechanical properties, and there is a problem in material stability. Furthermore, in a lightly machined portion represented by shallow drawing, since there is a high yield ratio, there is a problem that the springback becomes remarkable and the component accuracy cannot be increased.

  (3) Patent Document 3 discloses a thin steel sheet having a ferrite single phase structure in which fine precipitates having a particle size of less than 10 nm are dispersed and having a tensile strength of 550 MPa or more. By optimizing the hot rolling conditions, this steel sheet has a tensile strength of 700 MPa or more and a formability of a tensile strength-elongation balance (TS × El) of 17000 MPa ·% or more. However, considering the manufacturing process of cold-rolled steel sheets, which are thinner than hot-rolled steel sheets and have excellent surface roughness and thickness controllability, the addition of a large amount of carbonitride-forming elements increases the recrystallization temperature. In addition, since high temperature annealing is required, there is a problem that the coarsening of precipitates and the structure of the cold-rolled annealed steel sheet become coarse grains and the formability deteriorates.

  (4) Patent Document 4 discloses a low yield ratio high strength hot-rolled steel sheet using both precipitation strengthening and transformation strengthening. However, since a hard martensite phase is used, there is a problem that the bendability governed by the local ductility deteriorates as in the case of Patent Document 1. Furthermore, when precipitation strengthening is used in a hot-rolled steel sheet, it is difficult to control the amount of precipitates that control strength and the size of the precipitates, so it is easily inferred that there is a problem in material stability.

  (5) In Patent Document 5, in addition to 0.12 to 0.55 wt% C, 0.4 to 1.8 wt% Si, 0.2 to 2.5 wt% Mn, necessary After heating a steel sheet containing one or more of P, Ni, Cu, Cr, Ti, Nb, V, and Mo in an appropriate amount to the two-phase region of ferrite + austenite, A method for realizing a mixed structure of ferrite, bainite, and retained austenite by holding for 30 seconds to 30 minutes in a temperature range is disclosed.

  Patent Document 6 discloses a hot-rolled steel sheet containing 0.30 to 0.55% by weight of C, 0.7 to 2.0% by weight of Si, and 0.5 to 2.0% by weight of Mn. In addition, after heating the cold-rolled steel sheet to the austenite single-phase region, the steel plate is held at 650 to 750 ° C. for 4 to 15 seconds, and then held in the temperature range of 450 to 650 ° C. in the cooling process for a total of 10 to 50 seconds. Or the steel plate which improved the uniform deformability of the raw material using the processing induction transformation (TRIP) of the retained austenite which contains 10% or more of ferrite and 10% or more of retained austenite in the volume ratio in bainite is disclosed.

However, in any of these cases, the obtained steel sheet has a structure containing hard martensite after the processing-induced transformation, and thus there is a problem that the bendability governed by local ductility deteriorates. Furthermore, when considering the member performance after forming, sufficient impact characteristics cannot be obtained with members obtained after forming using these steel plates as raw materials.
JP-A-4-236671 JP-A-10-273754 JP 2002-322539 A Japanese Patent Laid-Open No. 5-179396 JP-A 61-157625 JP 60-43464 A

  A bag-like or cylindrical member formed by processing a steel plate is attached to an automobile after the seam is integrally joined by welding, for example, and the cross-section thereof is assembled into a closed cross-sectional shape. In the present specification, this is sometimes simply referred to as a closed section member. Therefore, the relationship between the deformation mode of the closed cross-section member during the actual collision and the material properties required for the steel sheet will be described.

FIG. 1 is a schematic view of a deformation mode of a cylindrical member having a closed cross-sectional shape used for a collision test.
When the collision direction in which a load is applied to the cylindrical member in the collision test is defined as the axial direction, the vertical cross section of the member with respect to the axial direction is configured by a plane portion between the ridge line portion having high rigidity and the ridge line portion having low rigidity (arc portion). ing. When a load is applied in the axial direction at the time of a collision, a flat portion having a low rigidity in the cross section causes elastic deflection in an outer direction of the cross section, and a compressive strain is generated in the axial direction in a ridge line portion having a high rigidity. Thereafter, as the crushing progresses, a difference occurs in the development behavior of the compressive strain in the thickness direction of the ridge line portion, and bending deformation at the ridge line portion, so-called plastic buckling (plastic deformation so that the ridge line portion is bent) occurs. Next, a wrinkle formed by buckling grows from the ridge line buckling start point to the flat surface portion, so that bending deformation also occurs in the flat surface portion. Thereafter, the wrinkles in the flat portion are crushed as the crushing progresses, and the wrinkles overlap and shift to the occurrence of plastic buckling in other ridge line portions.

  Thus, in the axial crushing deformation at the time of collision (plastic buckling deformation), plastic deformation including local bending deformation is repeated in the ridge line portion. This local bending deformation requires local ductility to withstand the deformation.

  In general, local ductility decreases with increasing strength. Therefore, when a member formed of a high-strength steel plate is subjected to deformation similar to that at the time of a collision, it cannot withstand local bending deformation, and a crack is likely to occur in the portion (buckling occurrence portion). The occurrence of such cracks causes a reduction in deformation load, and the desired energy absorption performance cannot be obtained.

In addition, when the closed cross-section is implemented by welding, even if the weld breaks, the deformation mode is disturbed and continuous plastic buckling cannot be realized, and the desired energy absorption capacity can be achieved. You will not be able to get.
On the other hand, if the work hardening characteristics decrease with increasing strength, the deformation mode will be disturbed, and it will not be possible to realize continuous fine plastic buckling in the impact test, and even if it does not lead to fracture, large bending deformation will occur. Occurs, the cylindrical member is greatly curved and deformed, and the desired energy absorption performance cannot be obtained. In the present specification, this is referred to as “disintegration” of the deformation mode at the time of collision. From the standpoint of absorbing collision energy, continuous plastic buckling "without collapse" is preferable.

  Therefore, in order to obtain a member having excellent collision performance, it has high local ductility and work hardening characteristics even in the member stage after molding, there is no crack at the buckling occurrence site, and there is no “collapse”. It is important to use as a raw material a steel plate that exhibits plastic buckling and has excellent resistance to fracture of welded parts (a characteristic that prevents the welded part constituting the closed section from breaking at the time of collision). . Fig. 4-1 (a), which will be described later, shows that the bent part was partially broken in the collision test. Figs. 4-2 (a) and (b) show that there is no "collapse". "Is seen in the schematic cross-sectional view of each side wall.

  Thus, high-strength steel sheets for satisfying the demands for promoting weight reduction of vehicles and improving collision safety include formability when forming into members, material stability that increases shape accuracy of members, and when assembling members. It is indispensable to obtain the desired stable impact absorption energy without causing breakage in the local bending deformation of the plastic buckling part at the time of impact deformation of the member and without causing breakage of the welded portion. .

  Therefore, the present invention is a problem of the conventionally proposed high-strength steel sheet, inferior formability due to inferior local ductility, lowering of the shape accuracy of members due to inferior material stability, bending deformation at the time of collision The purpose is to solve the breakage of welding, weldability at the time of assembly, and welded portion breakage at the time of collision.

 More specifically, an object of the present invention is to provide a high-strength steel sheet that can satisfy the formability at the time of forming the member and the collision property at the member stage after the forming process, and a method for manufacturing the same. .

  In the steel sheet according to the present invention, the target value of formability is a complex shape such as a pillar, which is considered to be a kind of automobile impact absorbing member, with a TS × El value obtained by a tensile test of 12000 MPa ·% or more. In order to cope with this, it is preferable that YR is 70% or less and the critical bending radius is 2.0 t or less.

As a result of repeated examinations of steel composition, steel structure, and production conditions, the present inventors have adjusted the steel composition and production conditions to an appropriate range, so that the area%, ferrite and bainite are 85% or more in total. Austenite is contained in an area ratio of 3.0 to 15%, the average particle size of the ferrite and bainite is 1.0 to 4.0 μm, and the C concentration in the retained austenite is 0.80 to 1.0% by mass; Furthermore, by making the steel structure containing 100 / μm ² or more of precipitates having a particle size of 1 to 10 nm in the ferrite and bainite, it is excellent in impact characteristics, material stability and formability while maintaining high strength. It was found that a high strength steel plate can be obtained.

  That is, it is necessary to improve the local deformability of the steel sheet in order to realize good formability to the member and to suppress cracking at the time of collision deformation at the member stage after forming. In order to improve local deformability, it is necessary to reduce the hardness difference between phases and structures constituting the steel structure and to improve the uniformity of the steel structure. However, the high-strength cold-rolled steel sheet that is the subject of the present invention cannot achieve the intended strength level by adopting a ferrite single-phase structure, so it must be made into a composite structure including other than the so-called ferrite phase such as bainite. I do not get.

  Therefore, it is necessary to obtain a mixed structure having a second phase that does not deteriorate the local deformability after having a uniform structure mainly composed of ferrite and bainite. The second phase that does not deteriorate the local deformability is a small amount of retained austenite phase, in which the C concentration in austenite is lower than the C concentration of austenite in ordinary retained austenitic steel, and is somewhat unstable as retained austenitic steel. It was found that retained austenite is effective.

  That is, at the time of molding, the moldability is improved while increasing the strength by the TRIP phenomenon due to the somewhat unstable retained austenite. Further, at the time of collision at the member stage after forming, the hardness of martensite generated from unstable retained austenite having a low C concentration is lower than the hardness of martensite generated from retained austenite of ordinary retained austenitic steel. Thus, a steel structure having a small hardness difference from the main phase can be realized, and the local ductility is excellent even in the member stage. Furthermore, by transforming from unstable retained austenite to martensite, the work hardening characteristics of the parts are also excellent.

  In addition, the hardness of such an unstable retained austenite having a low C concentration has a small strength fluctuation even if the amount of C in the steel fluctuates, so that the material stability can be improved compared to a steel sheet using martensite. it can.

  In this way, a high-strength steel sheet that can satisfy the formability at the time of forming the member and the collision property at the member stage after forming is obtained.

  According to the present invention, a high-strength cold-rolled steel sheet having an ultra-high strength with a tensile strength of 780 MPa or more and excellent in impact performance, material stability, and formability can be manufactured. The effect that contributes to the improvement of safety and collision safety is great.

Various conditions defined in the present invention will be described. First, the reason for limiting the steel composition will be described. In addition, unless otherwise indicated,% regarding steel composition means the mass%.
(C: 0.06-0.20%)
C is an austenite stabilizing element, which generates a retained austenite phase and effectively acts to increase the strength and improve the collision property. When Ti and Nb are contained as described later, soft ferrite and bainite are more easily generated during cooling than martensite. Therefore, in order to ensure a tensile strength of 780 MPa or more, at least 0.06% or more is contained. However, if the content exceeds 0.20%, the weldability deteriorates. For this reason, C amount was limited to 0.06 to 0.20% of range. In order to ensure a tensile strength of 980 MPa or more so that it can be applied to a single high-strength member, the content is preferably 0.10% or more.

(Si: 0.005-1.5%)
Si is an element that improves strength by solid solution strengthening and promotes ferrite transformation to improve ductility and the like. In the present invention, 0.005% or more is contained. However, if the content exceeds 1.5%, the nugget portion at the time of spot welding hardens and the toughness deteriorates. For this reason, the amount of Si was made into 0.005 to 1.5%. In addition, since chemical conversion processability may deteriorate when it contains more than 0.80%, it is preferably 0.80% or less.

(Mn: 1.6-3.0%)
Mn is an austenite stabilizing element and lowers the Ac ₃ transformation point. In order to facilitate the austenite single-phase region annealing during the continuous annealing, at least 1.6% or more is included. However, if the content exceeds 3.0%, soft ferrite cannot be obtained, so that ductility is deteriorated, and since a band-like structure is formed, bendability is deteriorated. For this reason, the amount of Mn was made 1.6 to 3.0%. In addition, if it is less than 1.8%, it will become difficult to make the annealing temperature used as an austenite single phase below 900 degreeC. On the other hand, if it exceeds 2.8%, the precipitation start temperature of ferrite becomes low, and it becomes hard ferrite and it is difficult to make YR 70% or less. For this reason, Preferably it is 1.8 to 2.8%.

(P: 0.03% or less)
P is an unavoidable impurity, and if it is contained excessively, it becomes a non-uniform band structure and the workability deteriorates. Therefore, the P content is set to 0.03% or less. In addition, Preferably it is 0.020% or less.

(S: 0.005% or less)
S exists as a sulfide in steel and is an element that greatly affects the improvement of local deformability targeted by the present invention. For this reason, it is desirable to reduce the amount of S as much as possible, but 0.005% or less If so, the characteristics as intended in the present invention are not adversely affected. In addition, Preferably it is 0.003% or less.

(Al: 0.3% or less)
Al is an element added to deoxidize steel and effectively acts to improve the yield of carbonitride-forming elements such as Ti. However, if the content exceeds 0.3%, the oxide inclusions increase, so the surface properties and moldability deteriorate. In addition, since the cost increases, the Al content is set to 0.3% or less. In addition, Preferably it is 0.02-0.08%.

(N: 0.01% or less)
N is an unavoidable impurity, and if it is contained excessively, coarse nitrides are precipitated, and formability deteriorates. For this reason, it is desirable to reduce the N amount as much as possible. However, if it is 0.01% or less, even a high-strength material as intended in the present invention does not adversely affect the moldability. Therefore, the N content is set to 0.01% or less. In addition, Preferably it is 0.005% or less, More preferably, it is 0.003% or less.

(Ti, Nb: 1 type or 2 types in total 0.03-0.25%)
These elements are one of the important elements in the present invention. In order to facilitate the formation of appropriate retained austenite by precipitation strengthening, strengthening the steel by forming carbides, grain refinement, and further promoting ferrite transformation. It works effectively. Moreover, there is an effect of making it difficult to cause HAZ softening when spot welding is performed.

  If the effect is less than 0.03%, the effect cannot be sufficiently obtained, so that not only ductility and collision characteristics deteriorate, but also YR cannot be made 70% or less. Therefore, at least 0.03% in total is contained. However, if the total content exceeds 0.25%, precipitates in the steel become coarse, and it becomes difficult to ensure a tensile strength of 780 MPa or more. If it is less than 0.05%, miniaturization becomes insufficient, so that the bendability is not sufficient, and if it exceeds 0.20%, it is difficult to make YR 70% or less. is there. For this reason, Preferably it is 0.05 to 0.20%.

(Cr, Mo, Ni: 0.5% or less each)
In addition to the above main elements serving as the skeleton of the present invention, one or more of Cr, Mo and Ni are appropriately selected within a range of 0.5% or less for the purpose of increasing strength and changing bainite transformation conditions. You may make it contain. In order to realize the increase in strength and the change in bainite transformation conditions more reliably, the content of elements to be contained is preferably 0.02% or more. On the other hand, when one or more of Cr, Mo, and Ni are contained in excess of 0.5%, ductility deteriorates because soft ferrite cannot be obtained.

The balance other than the components described above is Fe and impurities.
Next, the reason for limiting the steel structure of the high-strength cold-rolled steel sheet of the present invention will be described.
The high-strength cold-rolled steel sheet of the present invention having the above-described composition contains ferrite and bainite as the main phase, and in area ratio, contains a total of 85% or more of ferrite and bainite and 3.0 to 15% of retained austenite, The average particle size of the ferrite and bainite is 1.0 to 4.0 μm, the C concentration in the retained austenite is 0.80 to 1.0% by mass, and the particle size is 1 to 10 nm in the ferrite and bainite. Steel structure containing 100 precipitates / μm 2 or more.

(Area ratio of retained austenite: 3.0 to 15%, C concentration in austenite: 0.080 to 1.0% by mass)
In the present invention, the most important area ratio of retained austenite and C concentration in retained austenite are limited for the following reasons. If the retained austenite area ratio is less than 3.0%, the effect of improving the formability due to the TRIP effect is small, and if it exceeds 15%, the amount of martensite generated by the TRIP increases, and the microcrack occurs at the structure interface. Increases the local ductility, and the bendability at the time of forming the member and the impact characteristics at the member stage after forming are deteriorated.

  On the other hand, if the C concentration in the austenite is less than 0.80% by mass, the effect of improving the impact characteristics at the member stage after the forming process by the TRIP effect is small, and the C content dependency of the austenite hardness is increased, and the material stability is increased. Leading to a decline. On the other hand, if it exceeds 1.0% by mass, the strength of martensite alone generated by TRIP is increased, and microcracks are easily generated at the tissue interface, and the bendability at the time of forming the member and the member stage after the forming process. Adversely affects the collision characteristics.

For this reason, the area ratio of retained austenite is set to 3.0 to 15%, and the amount of C in the austenite is set to 0.80 to 1.0% by mass.
(Total area ratio of ferrite and bainite: 85% or more)
The steel structure of the steel sheet according to the present invention is a fraction evaluated by area ratio, and contains 85% or more of ferrite and bainite in total. When ferrite and bainite are included in a total area ratio of 85% or more, a desired retained austenite volume ratio and C content in retained austenite can be obtained, and a tensile strength of 780 MPa or more can be ensured without deteriorating collision characteristics and formability. It becomes possible.

(Average particle diameter of ferrite and bainite: 1.0 to 4.0 μm)
The steel structure of the steel sheet according to the present invention has an average particle size of ferrite and bainite of 1.0 to 4.0 μm. By setting the average particle size of ferrite and bainite to 4.0 μm or less, it is possible to ensure high strength without degrading formability. However, if the average particle size is less than 1.0 μm, the moldability deteriorates. For this reason, the average particle diameter of a ferrite and a bainite shall be 1.0-4.0 micrometers. In order to improve the bendability, it is preferable that the average particle size of ferrite and bainite is 3.2 μm or less.

  In order to make the average particle diameter of ferrite and bainite 1.0 to 4.0 μm, preferably 3.2 μm or less, it is preferable to contain an appropriate amount of alloy elements such as Ti, Nb, Mn, etc. as described above.

  In the present invention, the area ratio of ferrite and bainite, their average grain size, and the area ratio of retained austenite are within the scope of the present invention by adjusting, for example, the steel composition, hot / cold rolling conditions, and heat treatment conditions. The C content of retained austenite can also be adjusted, for example, by controlling the C content of the steel composition and the C diffusion amount during heat treatment.

(Precipitate density with particle size of 1 to 10 nm: 100 / μm ² or more)
Moreover, the structure of the steel sheet of the present invention contains precipitates having a particle size of 1 to 10 nm in the ferrite and bainite grains at a density of 100 / μm ² or more. Precipitates with a particle size greater than 10 nm do not act effectively on strengthening. On the other hand, if the number of precipitates having a particle size of 1 to 10 nm is less than 100 / μm ² , the amount of strengthening becomes small and the desired strength cannot be obtained. Precipitates at this time are Ti and / or Nb carbonitrides, which are precipitated by winding after hot rolling and heating and cooling during continuous annealing, and are hard and fine precipitates. Since it is a product, it greatly contributes to an increase in mechanical properties such as TS and YS.

In order to disperse precipitates having a particle diameter of 1 to 10 nm in the ferrite and bainite grains at a density of 100 pieces / μm ² or more, alloying elements such as Ti and Nb are contained in appropriate amounts, and hot rolling conditions described later. It is preferable to properly control the annealing conditions and the post-annealing cooling conditions. The upper limit of the density of the precipitate having a particle size of 1 to 10 nm is naturally limited by the equipment capacity, and is not particularly required to be defined within the range, but from the theoretical calculation, , Approximately 1000 / μm ² or less.

Next, the reason for limiting the manufacturing method of the high-strength cold-rolled steel sheet of the present invention will be described.
It is desirable that the molten steel having the above steel composition is melted by a generally known melting method such as a converter or an electric furnace, and used as a steel material such as a slab by a continuous casting method. In place of the continuous casting method, an ingot casting method, a thin slab casting method, or the like may be employed. This steel material is hot rolled to obtain a hot rolled steel sheet. In hot rolling, cast steel material is not cooled to room temperature but directly fed into a heating furnace while being heated and heated and then rolled, or directly after a little heat retention Rolling may be performed, or the steel material may be once cooled and then reheated for rolling.

(Reheating temperature of steel material: 1050-1300 ° C.)
In the case of reheating, it is necessary to re-dissolve TiC or NbC in order not to deteriorate the moldability. In the present invention, such an effect is recognized by heating to 1050 ° C. or higher, but heating to 1300 ° C. or higher not only saturates the effect but also increases scale loss. For this reason, the reheating temperature of the steel material at this time is preferably 1050 ° C to 1300 ° C. In other words, the hot rolling start temperature is preferably 1050 ° C to 1300 ° C.

  Moreover, in order to perform the said solid solution reliably, it is preferable to make heating time into 10 minutes or more, and in order to suppress an excessive scale loss, it is preferable to set it as 3 hours or less. More preferably, it is 30 minutes or more and 2 hours or less. Of course, when direct rolling or direct rolling is performed, as long as TiC and NbC are dissolved, rolling may be started as it is. In this case, the rolling start temperature is preferably 1050 to 1300 ° C.

(Finishing temperature: 800-950 ° C)
In this invention, finishing temperature shall be the range of 800 to 950 degreeC. If the finishing temperature is less than 800 ° C., the deformation resistance during rolling is large and the structure becomes a non-uniform band structure, and the formability of the steel sheet after continuous annealing deteriorates. On the other hand, when the temperature exceeds 950 ° C., grain growth occurs by subsequent cooling, and not only a uniform and fine structure cannot be obtained, but also precipitates become coarse and it becomes difficult to ensure strength after cold rolling and continuous annealing.

(Winding temperature: 450-750 ° C.)
In the present invention, the coiling temperature is in the range of 450 to 750 ° C. When the coiling temperature is less than 450 ° C., hard bainite and martensite are generated, and subsequent cold rolling becomes difficult. On the other hand, when the coiling temperature exceeds 750 ° C., the precipitate becomes coarse and it becomes difficult to ensure the strength after cold rolling and continuous annealing.

  The hot-rolled steel sheet is pickled by a normal method and then cold-rolled to obtain a cold-rolled steel sheet. In order to refine the structure of the steel sheet after the continuous annealing, it is preferable that the rolling reduction of the cold rolling is 30% or more.

(Heating temperature of cold-rolled steel sheet: above the temperature at which the cold-rolled steel sheet has an austenite single phase structure)
The cold-rolled steel sheet is annealed continuously and heated until the cold-rolled steel sheet has a temperature equal to or higher than an austenite single-phase structure (austenite single-phase temperature). Once the cold-rolled steel sheet has an austenite single-phase structure, the structure after continuous annealing is made uniform and fine. When the heating temperature is lower than the austenite single phase temperature, the concentration of C to austenite proceeds, the amount of C in the retained austenite contained in the steel sheet after continuous annealing increases, and not only the impact characteristics deteriorate, but also the cold-rolled structure The remaining effect becomes a band structure, and the bendability is remarkably deteriorated. In order to refine the precipitates contained in the steel sheet after continuous annealing and obtain a structure having a uniform crystal grain size, the average heating rate until the austenite single-phase structure is obtained is 1 ° C./second or more, 100 ° C./second. It is preferable to set it to 2 seconds or less.

(Annealing conditions of cold-rolled steel sheet: hold for 10 seconds or more in the austenite single phase state)
After heating to a range above the austenite single phase temperature, the austenite single phase structure is kept for 10 seconds or longer. When the holding time is less than 10 seconds, segregation of substitutional elements such as Mn remains, and the structure of the steel sheet after continuous annealing becomes non-uniform. For this reason, it was supposed that the annealing conditions of the cold-rolled steel sheet were maintained for 10 seconds or more in the austenite single phase state. In addition, since holding | maintenance of the austenite single phase structure | tissue for a long time raises the coarsening of an austenite particle size and it becomes difficult to make the structure | tissue after continuous annealing fine, it is preferable to make holding time into 300 second or less.

(Cooling conditions for cold-rolled steel sheet: The average cooling rate in the temperature range of 500 to 700 ° C is 200 ° C / second or less, and the cooling start temperature of ferrite is 600 to 750 ° C, and the temperature range of 150 to 300 ° C. To cool)
From the austenite single phase state, a cooling stop temperature range of 150 to 300 ° C. under cooling conditions in which the average cooling rate in the temperature range of 500 to 700 ° C. is 200 ° C./sec or less and the precipitation start of ferrite is 600 ° C. to 750 ° C. Allow to cool.

  When the average cooling rate of 500 to 700 ° C. exceeds 200 ° C./second, soft ferrite cannot be obtained, and not only ductility is deteriorated, but also ferrite does not precipitate, so C concentration in austenite cannot be performed. The collision characteristics deteriorate. For this reason, the average cooling rate in the temperature range of 500-700 degreeC shall be 200 degrees C / sec or less. If the average cooling rate is less than 5 ° C./second, strain cannot be imparted to the ferrite during cooling, and YR cannot be reduced to 70% or less. In addition, the cooling rate to 700 ° C. is not particularly limited as long as the ferrite does not precipitate above 750 ° C. after the austenite single-phase state.

  If the ferrite starts to be precipitated at a temperature exceeding 750 ° C., a soft ferrite phase is generated, so that a tensile strength of 780 MPa or more cannot be secured, and if it is less than 600 ° C., a hard ferrite phase is formed, and good ductility cannot be secured. For this reason, it is set as the cooling conditions from which the precipitation start temperature of a ferrite will be 600-750 degreeC. If ferrite precipitates at less than 630 ° C., it becomes hard ferrite and YR cannot be reduced to 70% or less. Therefore, the ferrite precipitation start temperature is preferably set to 630 ° C. or more.

  Although the cooling method is not specified, in order to further improve the flatness, material stability, and chemical conversion processability of the steel plate, avoid air / water cooling that can be rapidly cooled using water, avoid water cooling, and use gas cooling. Is preferred.

  In this invention, it cools to the cooling stop temperature range of 150-300 degreeC on the above cooling conditions. By controlling the cooling stop temperature to such a narrow range, the area ratio in the retained austenite and the amount of C in the retained austenite are set to desired values while having excellent impact characteristics while having a tensile strength of 780 MPa or more. In addition to facilitating the formation of martensite that leads to deterioration of material stability, the formation of bainite after cooling stop can be suppressed, and moreover, movable dislocations can be imparted to ferrite, and the yield ratio can be lowered. Preferably, the cooling stop temperature is 200 ° C. or higher.

(Condition after the cooling stop temperature: maintained in a temperature range of 150 to 300 ° C. for 30 to 1000 seconds)
After continuously cooling to a cooling stop temperature, the temperature is maintained in a temperature range of 150 to 300 ° C. for 30 to 1000 seconds, and then cooled to room temperature. In order to suppress the martensitic transformation and stabilize the austenite phase, the holding time at this time is 30 seconds or more. However, holding for 1000 seconds or more leads to waste of energy and a decrease in productivity.
Furthermore, it is preferable to perform temper rolling in the range of a rolling rate of 0.1 to 1%. Yield point elongation can be suppressed by temper rolling.

  Moreover, when corrosion resistance is calculated | required, you may give hot metal plating and electroplating to the steel plate surface. The plating type is not particularly limited as long as it is suitable for improving corrosion resistance. As the plating type, zinc or zinc alloy having sacrificial anticorrosive action is preferable, but other plating such as aluminum or aluminum alloy may be used.

Thus, by adjusting the steel composition, by optimizing continuous annealing conditions after hot rolling and cold rolling, the area ratio contains 85% or more of ferrite and bainite and 3 to 15% of retained austenite, The average particle size of the ferrite and bainite is 1.0 to 4.0 μm, the C concentration in the retained austenite is 0.80 to 1.0% by mass, and the particle size is 1 to 10 nm in the ferrite and bainite. A steel structure having 100 precipitates / μm ² or more can be obtained, and a high-strength cold-rolled steel sheet having a tensile strength of 780 MPa or more and excellent in formability, material stability, and impact characteristics can be obtained.

  Although the present invention has been described as a material for automobiles, the same characteristics are required not only for automobiles but also for various industrial machines. Therefore, as will be apparent to those skilled in the art from the above description, the present invention is applied. The material is not limited to automobile use.

  Here, in the present invention, `` formability derived from local ductility '' is the value of elongation (El) and tensile strength (TS) x elongation (El), `` shape accuracy of the member derived from material stability '' is Due to variations in tensile strength, “breakage at the bending deformation site at the time of collision” depends on whether or not there is a breakage in the collision test, “weldability at the time of assembly” is higher than the spot weldability test, and “ “Break” is also evaluated by the presence or absence of breakage in the collision test.

  Therefore, before describing the embodiment of the present invention, the collision test will be further described with reference to FIG.

Pre-crash test

  The present inventor press-molded a member as shown in FIG. 2 using a 980 MPa class high strength steel a having a chemical composition shown in Table 1 and a 590 MPa class high strength steel b as a comparison in order to investigate the impact performance. Welding the member to produce a member, and using a falling weight test facility, drop an 800 kg cone in the longitudinal direction of the member under the condition of a collision speed of 24 km / h, and investigate the deformation load at the time of collision and the form after deformation. did.

  The cylindrical member which is the specimen for the collision test in this example is shown in FIGS. 2 (a) and 2 (b) with a photograph and a schematic sectional view, respectively, and has a closed sectional shape by welding. The dimensions were as follows:

Board thickness: 1.2mm
Height: 170mm
Closed cross-sectional shape: recess-introduced rectangular tube shape The welding conditions were as follows.

Horizontal: Spot welding Electrode tip: 60φ, 40R, Chrome copper Power supply: AC Pressurization: 250kg
Current: 9kA
Current-carrying time: Fig. 3 shows the deformation load curve obtained in the 14-cycle test, and Fig. 4 shows the form photograph after deformation.

  As can be seen from FIG. 3, in the case of steel b, a flat deformation load that is comparatively stable in the entire deformation region is shown, whereas in the case of steel a, the deformation load is reduced after 60 mm. Also, from the deformed form of FIG. 4, steel b shows very fine buckling wrinkles and shows a stable deformed form, whereas the member of steel a can confirm the disorder of buckling wrinkles and break during deformation. However, it was found that stable plastic buckling could not be generated, and the generated buckling wrinkles were disturbed. Corresponding to high-speed video images, it was confirmed that rupture occurred during deformation, distortion of the deformation and instability of plastic buckling, and the cause of the decrease in deformation load was found.

  In this specification, the result of such a collision test is evaluated based on the presence or absence of fracture (including both fracture of the weld and fracture of the material) as the collision characteristics. The evaluation was based on “no collapse” and “collapse” corresponding to the case where the material was greatly bent and deformed.

  Fig. 4-1 is a photograph of the appearance of the molded product after the collision, Fig. 4-1 (a) is a partially enlarged side view of the steel a after the collision, and Fig. 4-1 (b) is a diagram of the steel a. Fig. 4-1 (c) shows the cross-section of the member after collision, and Fig. 4-1 (c) shows the cross-section of the post-impact member of steel b. Fig. 4-2 (a) shows plastic deformation so that the molded product is finely folded in the collision test. FIG. 4B is a schematic diagram of a cross section of the side wall portion when “no collapse”, and FIG. 4B is a schematic diagram of a cross section of the side wall portion when “no collapse”.

  When "collapse" occurs, as shown in the schematic diagram of Fig. 4-2 (b), the number of buckling wrinkles is small, and the member after the collision test is tilted from the axial direction. Can be determined. In fact, the member (test material) that was “collapsed” tilted in the direction of the arrow, and the number of buckling wrinkles was less than the member (test material) that was “no collapse”.

  A steel slab having the steel composition shown in Table 2 is heated to 1250 ° C., hot-rolled under the conditions shown in Table 3, and then wound, followed by cooling to a hot-rolled steel sheet (sheet thickness: 3.2 mm). ). Next, the hot-rolled steel sheet was pickled and cold-rolled to 1.0 mm to obtain a cold-rolled steel sheet. Thereafter, the cold-rolled steel sheet was heated to a temperature shown in Table 2 at a heating rate of 10 ° C./s, held for 60 s, and after annealing, gradually cooled to 660 ° C. at 3 ° C./s. It cooled at 60 degreeC / s to the cooling stop temperature shown, and hold | maintained 180 s at the said temperature.

  While confirming the austenite single phase formation in various heat treatment conditions of the steel pieces having the components shown in Table 2 and measuring the ferrite precipitation start temperature, the obtained cold-rolled annealed steel sheet was subjected to microstructure observation, X-ray residual austenite amount, austenite Using the measurement of C concentration in the inside, tensile test, and this material, the same member fabrication and falling weight test as in FIG. 2 were performed, and the forms of buckling wrinkles after impact deformation were compared.

Test methods for evaluating various production materials other than the collision test are shown below.
(experimental method)
(Confirmation of austenite single phase and measurement of ferrite precipitation start temperature)
Test specimens were collected from various cold-rolled steel sheets, and the change in expansion coefficient upon heat treatment as shown in Table 3 was analyzed to confirm the austenite single phase formation and the ferrite precipitation start temperature.

(Tissue observation)
Specimens were collected from the rolling direction of various cold-rolled annealed steel sheets and from the direction perpendicular to the rolling direction, and the cross-section in the rolling direction and the structure of the cross-section in the direction perpendicular to the rolling direction were photographed with an optical microscope or electron microscope. The fraction and the particle size of each phase were measured. The particle size was measured in accordance with the provisions of the cross line segment method of JISG 0552 for the total thickness of the sheet in the rolling direction cross section and the cross section perpendicular to the rolling direction, and the average value was expressed.

(Residual austenite amount and C concentration in austenite)
Various cold-rolled annealed steel sheets are subjected to chemical polishing to reduce the thickness by 0.3 mm, the surface after chemical polishing is subjected to X-ray diffraction, the amount of residual austenite is calculated, and the amount of austenite in the austenite is calculated based on the result. C concentration was calculated.

(Particle size and density of precipitates)
For the measurement of precipitate particle size and density, an electron microscope replica method was employed, five fields of view were photographed at a magnification of 100,000 for each sample, calculated as a circle-converted particle size, and precipitation with a particle size of 1 to 10 nm. The total number of objects was measured, and the number was divided by the area of the field of view to calculate the density.

(Tensile test)
JIS No. 5 tensile test specimens were taken from the direction perpendicular to the rolling direction of various cold-rolled annealed steel sheets and examined for yield strength (YS), tensile strength (TS), and elongation (El).

(Bending test)
A JIS No. 3 bending test piece having a longitudinal direction perpendicular to the rolling direction as a longitudinal direction was collected from various cold-rolled annealed plates, and the bendability was examined by a V-block method in accordance with JIS-Z-2248. At that time, a push fitting with an apex angle of 90 ° was pushed so that the burr was inside. The correctness after the test was visually inspected, and the minimum radius of the metal fitting with no crack observed after the test was divided by the plate thickness to calculate the minimum bending radius.

(Spot weldability test)
Spot weldability was performed under the conditions of a welding electrode tip diameter of 6 mm, an AC power source, a pressing force of 250 kg, a current of 9 kA, and an energization time of 14 cycles. After welding, the tensile load (TSS) by the tensile shear test of JISZ3136 and the tensile load (CTS) by the cross tensile test of JISZ3137 are measured. Those satisfying 0.35 or more were considered good.

(Explanation of test results)
These results are shown in Table 4. The steel sheet of the present invention contains, in area%, 85% or more of ferrite and bainite in total and 3.0 to 15% of retained austenite, and the average grain size of the ferrite and bainite is 1.0 to 4.0 μm. Yes, the C concentration in the retained austenite is 0.80 to 1.0% by mass, and further contains 100 precipitates / μm ² or more of precipitates having a particle size of 1 to 10 nm in the ferrite and bainite, and the tensile strength is It is a high-strength cold-rolled steel sheet having a formability of 780 MPa or more and a TS × El value of 12000 MPa ·% or more, good spot weldability, and the member after forming breaks or collapses during a collision test, and cracks in the welded part Does not occur.

  In contrast, the steel plate No. 1 and 2, the steel composition, that is, the chemical composition and the production conditions are out of the scope of the present invention, and not only the ductility is poor because the transformation start temperature of ferrite is lowered, but also the desired retained austenite volume fraction and retained austenite. Since the amount of C cannot be obtained, not only does the collapse occur during the collision test, but also the spot weldability is poor, and the weld fracture occurs during the collision test.

Steel plate No. In No. 7, the manufacturing conditions are out of the range of the present invention (specifically, the upper limit of the ferrite deposition start temperature is exceeded, and the ferrite becomes soft ferrite and the desired strength cannot be obtained.
Steel plate No. In No. 8, the chemical component is out of the scope of the present invention, and precipitates in the steel become coarse, and a desired strength cannot be obtained. Steel plate No. In No. 9, the chemical component is out of the scope of the present invention, and not only the desired strength cannot be obtained because bainite is soft, but also the amount of C in the retained austenite is low, and collapse occurs during a collision test. Steel plate No. Nos. 11 and 14 are out of the scope of the present invention, and most of the austenite is transformed into martensite, and the desired retained austenite volume fraction cannot be obtained or the amount of C in the retained austenite cannot be obtained. Therefore, collapse occurs during the collision test. Steel plate No. In No. 15, the production conditions deviate from the scope of the present invention, the concentration of C austenite progresses during annealing, the amount of C in the retained austenite increases, and cracks and collapse occur during a collision test. Steel plate No. In No. 17, the manufacturing conditions deviate from the scope of the present invention, and the transformation start temperature of ferrite becomes low, so that not only the ductility is bad, but also the amount of C in the retained austenite is low, and collapse occurs during the collision test. Steel plate No. In No. 18, the production conditions are out of the scope of the present invention, the cooling stop temperature and the subsequent holding temperature are high, ferrite and bainite are tempered, and the desired strength cannot be obtained. Steel plate No. No. 20 shows that the steel composition, that is, the chemical composition and production conditions are out of the scope of the present invention, and not only the ductility is poor due to the low transformation start temperature of ferrite, but also the desired retained austenite volume fraction and C in the retained austenite. Since the amount cannot be obtained, collapse occurs during the collision test.

  Among the examples of the present invention, steel plate No. 1 manufactured from steel type C having a component with a total content of Ti and Nb of less than 0.05%. No. 3 has an average particle size of ferrite and bainite of 3.2 μm or more and has poor bendability. In addition, among the inventive examples, the steel plate No. 1 manufactured from steel type D having a component with a C content of less than 0.10%. Nos. 4 and 5 are difficult to secure a higher tensile strength of 980 MPa or more. Moreover, among the examples of the present invention, the steel plate No. 1 manufactured from steel type E having a C content of less than 0.10% and a total content of Ti and Nb exceeding 0.20%. No. 6 is not only difficult to ensure a higher tensile strength of 980 MPa or more, but also does not have a YR of 70% or less. In addition, among the inventive examples, the steel plate No. manufactured under the condition that the average cooling rate from 700 ° C. to 500 ° C. is small. 12, YR does not become 70% or less. Moreover, among the examples of the present invention, a steel plate No. 1 manufactured from steel type K having a component having a Mn content of more than 2.8%. No. 19 has high hardenability, lowers the ferrite precipitation start temperature, and YR does not fall below 70%.

According to the steel sheet of the present invention example, it can be shown that variation in tensile strength can be suppressed.
The steel type I shown in Table 2 in the chemical composition range of the present invention example was used to investigate the dependence of the tensile strength on the cooling stop temperature. A cold-rolled sheet having a slab heating temperature of 1250 ° C., a finishing rolling temperature of 880 ° C., a finishing sheet thickness of 2.4 mm, a winding temperature of 600 ° C., and a cold rolling reduction ratio of 50% was used for the test. .

  Focusing on the cooling stop temperature, when manufacturing with a continuous annealing line as shown in FIG. 5, it is the last manufacturing condition to adjust the mechanical properties of the cold-rolled annealing plate, and the coil longitudinal direction and width direction For all of the above, the cooling stop temperature can only be controlled within 20 ° C. in the upper and lower directions, whereby the mechanical properties, in particular the tensile strength in question, tend to vary.

  FIG. 6 shows the cooling stop temperature dependence of the tensile strength obtained in the test. The results showed that the dependence of the tensile strength on the cooling stop temperature was small at 300 ° C. or lower and increased when the temperature exceeded 300 ° C. Further, as described above, since the cooling stop temperature can be controlled only within 20 ° C. above and below the target temperature, the variation in tensile strength may be considered as the maximum difference in tensile strength between the target temperature and 20 ° C. above and below. Therefore, the tensile strength was measured when 20 ° C. above and below the target temperature was the cooling stop temperature, and the difference was taken as the strength variation. Specifically, the strength variation at 300 ° C. is a difference in tensile strength when 280 ° C. and 320 ° C. are set as the cooling stop temperatures.

  As a result of analyzing the dependence of the tensile strength on the cooling stop temperature, when the temperature is 300 ° C. or less, the strength variation is suppressed within 5% of the desired tensile strength, whereas when the temperature exceeds 300 ° C., the variation is 5% or more. Was encouraged. On the other hand, when the temperature exceeds 400 ° C., although the variation in tensile strength is being suppressed again, a tensile strength of 780 MPa or more cannot be obtained. Therefore, it was found that the cooling stop temperature must be 300 ° C. or lower in order to suppress the variation in tensile strength within 5% while securing the tensile strength of 780 MPa or more.

FIGS. 1A and 1B are schematic diagrams of member deformation at the time of collision when an impact load is applied to a cylindrical member which is a test material in a collision test, FIG. 1A is just after the collision, FIG. (c) is a schematic diagram of each after the collision. FIG. 2A is a photograph showing a side surface of the simulated molded product used in the collision test of FIG. 1, and FIG. 2B is a schematic diagram of the horizontal section thereof. The load displacement curve at the time of a collision test is shown. Fig. 4-1 is a photograph of the appearance of the molded product after the collision, Fig. 4-1 (a) is a partially enlarged side view of the steel a after the collision, and Fig. 4-1 (b) is a diagram of the steel a. Fig. 4-1 (c) shows the cross-section of the member after collision, and Fig. 4-1 (c) shows the cross-section of the post-impact member of steel b. Fig. 4-2 (a) shows plastic deformation so that the molded product is finely folded in the collision test. FIG. 4B is a schematic diagram of a cross section of the side wall portion when “no collapse”, and FIG. 4B is a schematic diagram of a cross section of the side wall portion when “no collapse”. It is a temperature change figure which shows continuous annealing conditions. It is a graph of the test result which shows the cooling stop temperature dependence of tensile strength.

Claims (3)

In mass%, C: 0.06-0.20%, Si: 0.005-1.5%, Mn: 1.6-3.0%, P: 0.03% or less, S: 0.005 % Or less, Al: 0.3% or less, N: 0.01% or less, and one or two of Ti and Nb are contained in a total of 0.03 to 0.25%, with the balance being Fe and impurities. The steel composition has a total area of 85% or more of ferrite and bainite and 3.0 to 15% of retained austenite, and the average grain size of the ferrite and bainite is 1.0 to 4.0 μm. A steel structure containing 100 / μm ² or more of precipitates having a particle size of 1 to 10 nm in the ferrite and bainite, the C concentration in the retained austenite is 0.80 to 1.0% by mass, and tensile The strength is 780 MPa or more. Cold-rolled steel sheet to be.   The steel composition is one or more selected from the group consisting of Cr: 0.5% or less, Mo: 0.5% or less, and Ni: 0.5% or less, instead of part of the Fe. The cold-rolled steel sheet according to claim 1, comprising: The method for producing a high-strength cold-rolled steel sheet according to claim 1 or 2, comprising the following steps (A) to (C):
(A) Hot rolling a hot rolled steel sheet by subjecting a steel material having the steel composition according to claim 1 or 2 to hot rolling at a finishing temperature of 800 ° C. to 950 ° C. and a winding temperature of 450 to 750 ° C. Process;
(B) a cold rolling step in which the hot-rolled steel sheet is cold-rolled to obtain a cold-rolled steel sheet; and (C) the cold-rolled steel sheet is held in an austenite single phase structure for 10 seconds or more, and then 500 It is cooled to a cooling stop temperature range of 150 to 300 ° C. under cooling conditions in which an average cooling rate in a temperature range of ˜700 ° C. is 200 ° C./second or less and a ferrite precipitation start temperature is 600 to 750 ° C. A continuous annealing step in which a heat treatment is performed for 30 to 1000 seconds in the temperature range.

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