JP2015193897A - High strength cold rolled steel sheet and high strength hot-dip galvanized steel sheet having excellent ductility and bendability, and methods for producing the same - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high strength cold rolled steel sheet which has a tensile strength of 980 MPa or greater and which exhibits excellent ductility and bendability.SOLUTION: There is provided the high strength cold rolled steel sheet which has a tensile strength of 980 MPa or greater and exhibits excellent ductility and bendability, and satisfies a specified component composition, and in which structure at a depth corresponding to a quarter of the sheet thickness satisfies all of (1) to (5) below when measured using specified methods. (1) The areal ratio of ferrite relative to the overall structure is 5% or more but less than 50%, with the remainder being hard phases. (2) The areal ratio of a mixed structure of fresh martensite and residual austenite relative to the overall structure is greater than 0% but 30% or less (3) A region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is present at 5% by area or greater. (4) When the fraction of a region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is measured for 100 sections in a section of 2 μm square, the standard deviation is 4.0% or greater. (5) The concentration of Mn in a ferrite phase is 0.9 times or less the concentration of Mn in the steel sheet.

Description

  The present invention relates to a high-strength cold-rolled steel sheet and a high-strength hot-dip galvanized steel sheet excellent in ductility and bendability, and methods for producing them. Specifically, in a region where the tensile strength is 980 MPa or more, a high-strength cold-rolled steel sheet having high ductility and bendability, a high-strength electrogalvanized steel sheet, a high-strength hot-dip galvanized steel sheet, and a high-strength galvannealed steel sheet, The present invention relates to a production method capable of producing these steel plates efficiently.

  In order to reduce fuel consumption of automobiles and transport aircraft, it is desired to reduce the weight of automobiles and transport aircraft. For example, it is effective to use a high-strength steel plate and reduce the thickness to reduce the weight. However, when the strength of the steel plate is increased, the ductility deteriorates and the workability deteriorates. Therefore, high strength steel sheets are required to have excellent ductility and bendability necessary for press forming. In addition, from the viewpoint of corrosion resistance, steel parts for automobiles are galvanized steel sheets such as electrogalvanized (EG), hot dip galvanized (GI), and alloyed hot dip galvanized (GA). Is often used as a representative). These galvanized steel sheets are also required to have the same characteristics as high-strength steel sheets.

  As techniques for improving the workability of high-strength steel sheets, for example, Patent Documents 1 to 4 propose ultra-high-strength cold-rolled steel sheets with excellent bendability.

JP 2011-179030 A Japanese Patent No. 5299591 JP 2011-225976 A International Publication No. 2012/036269

  However, a steel sheet excellent in ductility and bendability in a region where the tensile strength is 980 MPa or more has not yet been provided.

  The present invention has been made paying attention to the above-mentioned circumstances, and the purpose thereof is a high-strength cold-rolled steel sheet and a high-strength hot-dip galvanized plate having a tensile strength of 980 MPa or more and excellent ductility and bendability. Steel sheets and methods for producing them with high productivity, specifically, high strength cold-rolled steel sheets, high-strength electrogalvanized steel sheets, high-strength hot-dip galvanized steels with excellent ductility and bendability in the region where the tensile strength is 980 MPa or more An object of the present invention is to provide a steel plate, a high-strength galvannealed steel plate, and a production method capable of efficiently producing these steel plates.

The present invention that can solve the above-mentioned problems is that the composition of the steel sheet is, by mass, C: 0.10% to 0.30%, Si: 1.2% to 3%, Mn: 0.00. 5% to 3.0%, P: more than 0% to 0.1%, S: more than 0% to 0.05%, Al: 0.005% to 0.2%, N: more than 0% 0 .01% or less, and O: more than 0% and 0.01% or less are satisfied, the balance is made of iron and unavoidable impurities, and the structure of the steel plate at the 1/4 position is the following (1) to (5 ) To meet all of the above. Hereinafter, “%” means “mass%” for the component composition of the steel sheet.
(1) When observed with a scanning electron microscope, the area ratio of ferrite to the entire structure is 5% or more and less than 50%, and the remainder is a hard phase.
(2) When the repeller corrosion is performed and observed with an optical microscope, the area ratio of the mixed structure of fresh martensite and retained austenite with respect to the entire structure is more than 0% and 30% or less.
(3) When analyzed with an electron beam microprobe analyzer, there are 5 area% or more of regions where the Mn concentration is 1.2 times or more the Mn concentration in the steel sheet, and (4) 2 μm sections The fraction of the region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel sheet is measured, and the standard deviation when measuring 100 sections is 4.0% or more.
(5) When analyzed with an electron microprobe analyzer, the Mn concentration in the ferrite phase is 0.90 times or less of the Mn concentration in the steel sheet.

  In the present invention, it is also a preferred embodiment that the volume ratio of retained austenite with respect to the entire structure is 5% or more when measured by the X-ray diffraction method.

  The hard phase preferably comprises a mixed structure of the fresh martensite and retained austenite; and at least one structure selected from the group consisting of bainitic ferrite, bainite, and tempered martensite.

  In practicing the present invention, as another element, (A) Cr: at least one selected from the group consisting of more than 0% and not more than 1% and Mo: more than 0% and not more than 1%; (B) Ti: 0 %: At least one selected from the group consisting of more than 0.15% or less, Nb: more than 0% and 0.15% or less, and V: more than 0% and less than 0.15%; (C) Cu: more than 0% and 1% And at least one selected from the group consisting of Ni: more than 0% and 1% or less; (D) B: more than 0% and 0.005% or less; (E) Ca: more than 0% and 0.01% or less; Mg It is also a preferred embodiment that contains at least one selected from the group consisting of: more than 0% and not more than 0.01%, and REM: more than 0% and not more than 0.01%.

  The present invention further includes a high-strength electrogalvanized steel sheet having an electrogalvanized layer formed on the surface of the high-strength cold-rolled steel sheet; a hot-dip galvanized layer formed on the surface of the high-strength cold-rolled steel sheet. And a high-strength galvanized steel sheet having an alloyed galvanized layer formed on the surface of the high-strength cold-rolled steel sheet.

The method for producing a high-strength cold-rolled steel sheet according to the present invention that has solved the above problems is a hot-rolling step of a steel sheet having the above component composition, and is wound at a winding temperature of 500 ° C. or more and 800 ° C. or less, and then 500 ° C. Hold at 800 ° C. or lower for 3 hours or longer, then cool to room temperature, and after cold rolling, keep soaked in a temperature range of (Ac ₁ point + 20 ° C.) to less than Ac ₃ point, and then average cooling rate to 500 ° C. Cool from 10 ° C / second to 500 ° C at an average cooling rate of 10 ° C / second to a temperature range of 500 ° C or lower, then reheat to a temperature range of 250 ° C to 500 ° C and hold for 30 seconds or longer Then, it has a gist in cooling to room temperature.

  In the present invention, it is also preferred that the steel sheet obtained by the above production method is further subjected to electrogalvanization.

Moreover, the manufacturing method of the high intensity | strength hot-dip galvanized steel plate which concerns on this invention which could solve the said subject is a hot-rolling process of the steel plate which consists of the said component composition, and is wound up by coiling temperature 500 degreeC or more and 800 degrees C or less, and then 500 Hold at 3 ° C to 800 ° C for 3 hours or more, then cool to room temperature, and after cold rolling, keep soaked in a temperature range of (Ac ₁ point + 20 ° C) to less than Ac ₃ point, then average cooling to 500 ° C Cool at a rate of 10 ° C./second or more and 500 ° C. or less at an average cooling rate of 10 ° C./second or more to a temperature range of 500 ° C. or less, and then reheat to a temperature range of 250 ° C. or more and 500 ° C. or less, for 30 seconds or more The main point is to hold and to cool to room temperature after hot dip galvanization within the holding time.

  In the present invention, it is also preferable to perform alloying in a temperature range of 450 ° C. or more and 550 ° C. or less after performing the hot dip galvanization.

  According to the present invention, a high-strength cold-rolled steel sheet and a high-strength hot-dip galvanized steel sheet that are excellent in ductility and bendability even at 980 MPa or more, in particular, a high-strength cold-rolled steel sheet that has excellent characteristics, a high strength An electrogalvanized steel sheet, a high-strength hot-dip galvanized steel sheet, and a high-strength galvannealed steel sheet can be provided. Moreover, according to the manufacturing method of this invention, these steel plates can be manufactured efficiently. Therefore, the high-strength cold-rolled steel sheet of the present invention is extremely useful particularly in the industrial field such as automobiles.

FIG. 1 is a schematic explanatory view showing an example of a heat treatment pattern in the production method of the present invention.

  In order to improve the ductility and bendability of particularly high-strength cold-rolled steel sheets and high-strength hot-dip galvanized steel sheets having a tensile strength of 980 MPa or more, the present inventors have made extensive studies.

  As a result, assuming that the component composition is appropriately controlled, the ferrite phase and the hard phase in the metallographic structure of the steel sheet are optimized, and if segregation of Mn is appropriately controlled, a high strength of 980 MPa or more is secured. The present inventors have found that ductility and bendability can be improved, and have reached the present invention.

  Hereinafter, the reason for defining the metal structure in the present invention will be described in detail. In addition, the fraction measured by microscopic observation means the ratio for the whole structure (100%) of a steel plate. The measurement method of the metal structure constituting the present invention differs depending on the metal structure. Therefore, when all the metal structures defined in the present invention are totaled, it may exceed 100%. This is because the residual γ constituting the mixed structure of fresh martensite and residual austenite is only measured by optical microscope observation. This is because it is also measured by X-ray diffraction. Hereinafter, retained austenite is referred to as “residual γ”, and the mixed structure of fresh martensite and retained austenite is sometimes referred to as “MA (Martensite-Austenite Constituent) structure”.

[Area ratio of ferrite: 5% or more and less than 50%]
Ferrite is a structure having an effect of improving the ductility and bendability of a steel sheet. In the present invention, ductility and bendability in a high strength region having a tensile strength of 980 MPa or more can be improved by increasing the area fraction of ferrite. In order to exert such an effect, the area ratio of ferrite is 5% or more, preferably 7% or more, more preferably 10% or more. However, when the ferrite is excessive, the strength of the steel sheet is lowered, and it becomes difficult to ensure a high strength of 980 MPa or more. Therefore, the area ratio of ferrite is less than 50%, preferably 45% or less, more preferably 40% or less. The area ratio of the ferrite is a value obtained by observing the position of the steel sheet with a thickness of 1/4 by observation with a scanning electron microscope (SEM: Scanning Electron Microscope).

[Hard phase]
The hard phase is a structure necessary for improving the tensile strength. In the present invention, by increasing the area fraction of the hard phase, high strength of 980 MPa or more can be achieved while allowing soft ferrite to exist within the range of the area ratio. In order to exert such an effect, the remaining metal structure other than ferrite needs to be a hard phase. In the present invention, the hard phase is a phase harder than ferrite, and is at least one selected from the group consisting of bainitic ferrite, bainite, tempered martensite, and MA structure. In the present invention, the following is shown. As such, it includes at least the MA organization. Among the hard phases, bainitic ferrite, bainite, and tempered martensite are measured values by SEM observation at a 1/4 thickness position of the steel sheet. Residual γ is included between the laths of bainitic ferrite or in the MA structure.

[Area ratio of MA structure to all structures: more than 0% and 30% or less]
If the MA structure is present, the strength and ductility can be improved. Therefore, from the viewpoint of improving the strength-ductility balance, the area ratio of the MA structure is preferably 3% or more, more preferably 4% or more. On the other hand, if the area ratio of the MA structure is too large, the bendability deteriorates. Therefore, in the present invention, the area ratio of the MA structure is 30% or less, preferably 20% or less, more preferably 15% or less.

  In addition, the fresh martensite which comprises MA structure means the state in which the untransformed austenite was martensitic transformed in the process of cooling a steel plate from heating temperature to room temperature, and is distinguished from tempered martensite after heat treatment. . In the present invention, the portion that has been whitened when observed with an optical microscope due to repeller corrosion is defined as the MA structure. Since fresh martensite and residual γ are difficult to distinguish by observation with an optical microscope, a composite structure of fresh martensite and residual γ is measured as an MA structure. The MA structure is a value measured by observation with an optical microscope at a ¼ position of the steel sheet thickness.

[A region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel plate: 5 area% or more, and a region where the Mn concentration is concentrated 1.2 times or more the Mn concentration in the steel plate in □ 2 μm section. Standard deviation of the area fraction of the product: 4.0% or more]
In the present invention, the region where the Mn concentration is concentrated is the Mn obtained by analyzing the cross section of the steel sheet with a beam diameter of 1 μm or less and 20 μm × 20 μm using an electron probe microanalyzer (EPMA). It is defined using the concentration. The “Mn concentration in the steel plate” is a Mn concentration obtained by chemical analysis of the base steel plate by inductively coupled plasma emission spectroscopy. Therefore, the region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel sheet means that the measured value of the Mn concentration obtained by EPMA analysis is 1.2 times or more higher than the Mn concentration in the base steel sheet. This is a region and is measured in the range of 20 μm × 20 μm.

  The “□ 2 μm section” is a 2 μm square section. In the present invention, an EPMA measurement range of 20 μm × 20 μm is divided into 100 sections of 1 μm 2 μm square obtained by drawing lines 2 μm apart in length and width. In each section, the area fraction of the region where the Mn concentration is 1.2 times or more is measured, and the standard deviation is statistically obtained in 100 sections.

  In the present invention, in the Mn concentration distribution, a region where the concentration of Mn in the steel plate is concentrated 1.2 times or more is 5 area% or more, and Mn is concentrated 1.2 times or more in the square of 2 μm. It has been found that if the standard knitting difference when measuring the area fraction is 4.0% or more, the bendability is greatly improved. Below, the region where the Mn concentration in the steel sheet is concentrated 1.2 times or more is referred to as the “region where the Mn concentration is 1.2 times or more”, and Mn is concentrated 1.2 times or more in the 2 μm section. The standard stitch difference when the area fraction is measured may be referred to as “standard deviation of a region where the Mn concentration is concentrated 1.2 times or more” or simply “standard deviation”.

  That is, the region having a Mn concentration of 1.2 times or more mainly becomes a hard phase. And, as the area ratio of the region having a Mn concentration of 1.2 times or more is larger, the Mn concentration in the ferrite phase is relatively lowered, the hardness of the ferrite phase can be lowered, and the bendability can be improved. The standard deviation increases as the amount of segregation of Mn increases. However, the Mn concentration in the ferrite phase that contributes to improvement in bendability decreases, and the hardness of the ferrite phase can be reduced.

  In order to obtain such an effect, the region having a Mn concentration of 1.2 times or more is 5.0 area% or more, preferably 5.2 area% or more, more preferably 5.5 area% or more. On the other hand, if the ratio of the Mn concentration of 1.2 times or more is too high, the Ms point of austenite may decrease and the MA structure may increase, so the region of Mn concentration of 1.2 times or more is preferably 20 areas. % Or less, more preferably 15 area% or less.

  The standard deviation of the region where the Mn concentration is 1.2 times or more is 4.0% or more, preferably 4.5% or more, more preferably 5.0% or more. When the standard deviation is smaller than 4.0%, since the distribution of Mn is insufficient and uniform, the hardness reduction of the ferrite phase that contributes to improvement of bendability is insufficient. On the other hand, the upper limit of the standard deviation is not particularly limited, and is preferably 10% or less.

  And by making the region of Mn concentration 1.2 times or more 5 area% or more and the standard deviation 4.0% or more, the spheroidized hard phase can be dispersed in the ferrite phase, and the strength of the steel material The improvement effect and the bendability improvement effect by ferrite can be combined. Hereinafter, the spheroidized hard phase is referred to as “spherical hard phase”. Here, the spherical hard phase is a part of the hard phase, and is composed of bainitic ferrite, bainite, tempered martensite, MA structure and the like as the hard phase. Conventionally, it was thought that when the hardness difference between the hard phase and the ferrite phase was large, an interface crack was generated during bending and the bendability was deteriorated. However, it was found that the interface crack can be suppressed by the spherical hard phase in the ferrite phase. In order to obtain such an effect, the spherical hard phase in the ferrite phase should be small, and the aspect ratio is preferably 3 or less, more preferably 2.5 or less, still more preferably 2 or less, and the equivalent circle diameter. Is preferably 2 μm or less, more preferably 1.8 μm or less, and still more preferably 1.5 μm or less. In order to achieve the above effect, the spherical hard phase is preferably 0.70% by volume or more, more preferably 0.75% by volume or more, and further preferably 0.80% by volume or more with respect to the hard phase. To do.

[Mn concentration in ferrite is less than 0.90 times Mn concentration in steel sheet]
If the Mn concentration in the ferrite phase is too high, the hardness of the ferrite phase cannot be sufficiently reduced, and the bendability deteriorates. Therefore, the Mn concentration in the ferrite phase needs to be lower than the Mn concentration in the steel sheet. Therefore, the Mn concentration in the ferrite phase is 0.90 times or less, preferably 0.85 times or less, more preferably 0.80 times or less the Mn concentration in the steel sheet. On the other hand, if the Mn concentration in the ferrite is too low, the ferrite hardness decreases and the strength may be insufficient. Therefore, the Mn concentration in the ferrite is preferably 0.3 times or more the Mn concentration in the steel sheet. Preferably it is 0.4 times or more. The Mn concentration in the ferrite phase can be measured by EPMA.

[Volume ratio of residual γ with respect to all tissues: 5% or more]
Residual γ is deformed in response to the deformation of the steel sheet and transforms into martensite, thereby ensuring good ductility and promoting the hardening of the deformed part during processing, thereby suppressing the concentration of strain. Therefore, it is a structure necessary for improving the strength-ductility balance of the steel sheet. In order to effectively exhibit such an effect, the volume ratio of the residual γ is preferably 5% or more, more preferably 6% or more, and further preferably 7% or more. The upper limit of the volume ratio of residual γ is not particularly limited, but is 20% or less at most within the range of the component composition and production conditions of the present invention. Residual γ is a value measured by an X-ray diffraction method at a position of 1/4 of the thickness of the steel sheet.

  Residual γ is present between the laths of bainitic ferrite or in the MA structure. Since the effect of the residual γ is exhibited regardless of the existence form, in the present invention, the residual γ that can be confirmed upon measurement is the residual γ regardless of the existence form.

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

[C: 0.10% to 0.30%]
C is an element necessary for ensuring strength and enhancing the stability of residual γ. In order to ensure a tensile strength of 980 MPa or more, the C content is 0.10% or more, preferably 0.12% or more, more preferably 0.15% or more. However, if the C content is excessive, the strength after hot rolling increases, cracking occurs during cold rolling, and the weldability of the final product decreases, so the C content is 0.30% or less, preferably It is made 0.26% or less, more preferably 0.23% or less.

[Si: 1.2% to 3%]
Si is an element that contributes to increasing the strength of steel as a solid solution strengthening element. Moreover, it is an element effective in suppressing the generation of carbides, effectively acting on the generation of residual γ, and ensuring an excellent TS × EL balance. In order to effectively exert such effects, the Si content is set to 1.2% or more, preferably 1.35% or more, more preferably 1.5% or more. However, when the Si content is excessive, a significant scale is formed during hot rolling, and scale marks are formed on the surface of the steel sheet, which may deteriorate the surface properties. Moreover, pickling property is deteriorated. Therefore, the Si content is 3% or less, preferably 2.8% or less, more preferably 2.6% or less.

[Mn: 0.5% to 3.0%]
Mn is an element that contributes to increasing the strength of the steel sheet by improving the hardenability. Further, it is an element that effectively acts to stabilize γ and generate residual γ. In order to effectively exert such action, the Mn content is 0.5% or more, preferably 0.6% or more, more preferably 1.0% or more, still more preferably 1.5% or more, and still more. Preferably it is 2.0% or more. However, if the Mn content is excessive, the strength after hot rolling is increased, causing cracks during cold rolling, or causing the weldability of the final product to deteriorate. Excessive Mn addition causes segregation of Mn and deteriorates workability. Therefore, the Mn content is 3.0% or less, preferably 2.8% or less, more preferably 2.6% or less.

[P: more than 0% and 0.1% or less]
P is an element unavoidably contained, and is an element that deteriorates the weldability of the steel sheet. Therefore, the P content is 0.1% or less, preferably 0.08% or less, more preferably 0.05% or less. In addition, since it is better that the P content is as small as possible, the lower limit is not particularly limited, but industrially the lower limit is 0.0005%.

[S: more than 0% and 0.05% or less]
S, like P, is an element that is inevitably contained, and is an element that degrades the weldability of the steel sheet. Further, S forms sulfide inclusions in the steel sheet and causes the workability of the steel sheet to deteriorate. Therefore, the S content is 0.05% or less, preferably 0.01% or less, more preferably 0.005% or less. Since the S content is preferably as low as possible, the lower limit is not particularly limited, but industrially the lower limit is 0.0001%.

[Al: 0.005% to 0.2%]
Al is an element that acts as a deoxidizer. In order to effectively exhibit such an action, the Al content is set to 0.005% or more, more preferably 0.01% or more. However, if the Al content is excessive, the weldability of the steel sheet is remarkably deteriorated, so the Al content is 0.2% or less, preferably 0.15% or less, more preferably 0.10% or less.

[N: more than 0% and 0.01% or less]
N is an element inevitably contained, but is an element that contributes to increasing the strength of the steel sheet by precipitating nitrides in the steel sheet. In this respect, the N content is preferably 0.001% or more. However, if the N content is excessive, a large amount of nitride precipitates and stretches, causing deterioration of stretch flangeability (λ), bendability, and the like. Therefore, the N content is 0.01% or less, preferably 0.008% or less, more preferably 0.005% or less.

[O: more than 0% and 0.01% or less]
O is an element that is unavoidably included, and if excessively included, it is an element that causes a decrease in ductility and bendability during processing. Therefore, the O content is 0.01% or less, preferably 0.005% or less, more preferably 0.003% or less. In addition, since it is better that the O content is as small as possible, the lower limit is not particularly limited, but industrially the lower limit is 0.0001%.

[Other ingredients]
The steel sheet of the present invention satisfies the above component composition, and the balance is iron and inevitable impurities. The inevitable impurities include, for example, the above-mentioned P, S, N, O, and other trump elements such as Pb, Bi, Sb, Sn that may be brought into steel depending on the conditions of raw materials, materials, manufacturing equipment, and the like. Sometimes. Moreover, it is also possible to positively contain the following elements as other elements as long as the effects of the present invention are not adversely affected.

The steel sheet of the present invention is further as another element,
(A) at least one selected from the group consisting of Cr: more than 0% and 1% or less and Mo: more than 0% and 1% or less,
(B) Ti: at least one selected from the group consisting of more than 0% and 0.15% or less, Nb: more than 0% and 0.15% or less, and V: more than 0% and 0.15% or less,
(C) at least one selected from the group consisting of Cu: more than 0% and 1% or less and Ni: more than 0% and 1% or less,
(D) B: more than 0% and 0.005% or less,
(E) Ca: more than 0% 0.01% or less, Mg: more than 0% 0.01% or less, and REM: at least one selected from the group consisting of more than 0% and 0.01% or less. May be. These elements (A) to (E) may be contained alone or in any combination. The reason for setting this range is as follows.

[(A) Cr: at least one selected from the group consisting of more than 0% and 1% or less and Mo: more than 0% and 1% or less]
Cr and Mo are both effective elements for improving the hardenability and improving the strength of the steel sheet, and can be used alone or in combination. In order to effectively exhibit such an action, the contents of Cr and Mo are preferably 0.1% or more, more preferably 0.3% or more, respectively. However, if it is excessively contained, the workability is reduced and the cost is increased. Therefore, when Cr and Mo are contained alone, preferably 1% or less, more preferably 0.8% or less, More preferably, it is 0.5% or less. When Cr and Mo are used in combination, each is independently within the above upper limit range, and preferably the total amount is 1.5% or less.

[(B) At least one selected from the group consisting of Ti: more than 0% and 0.15% or less, Nb: more than 0% and 0.15% or less, and V: more than 0% and 0.15% or less]
Ti, Nb, and V are all elements that have the effect of forming carbide and nitride precipitates in the steel sheet, improving the strength of the steel sheet, and refining the old γ grains, alone or Can be used in combination. In order to effectively exhibit such an action, the contents of Ti, Nb, and V are each preferably 0.005% or more, more preferably 0.010% or more. However, when it contains excessively, carbide will precipitate to a grain boundary and the stretch flangeability and bendability of a steel plate will deteriorate. Therefore, the contents of Ti, Nb and V are each preferably 0.15% or less, more preferably 0.12% or less, and still more preferably 0.10% or less.

[(C) Cu: at least one selected from the group consisting of more than 0% and 1% or less and Ni: more than 0% and 1% or less]
Cu and Ni are elements that effectively act to generate and stabilize retained austenite, and also have an effect of improving corrosion resistance, and can be used alone or in combination. In order to exert such an effect, the contents of Cu and Ni are preferably 0.05% or more, more preferably 0.10% or more, respectively. However, since hot workability deteriorates when Cu is contained excessively, the Cu content is preferably 1% or less, more preferably 0.8% or less, and even more preferably 0.5% when added alone. % Or less. When Ni is excessively contained, the cost becomes high, so the Ni content is preferably 1% or less, more preferably 0.8% or less, and still more preferably 0.5% or less. When Cu and Ni are used in combination, the above-described effect is easily exhibited, and the deterioration of hot workability due to the addition of Cu is suppressed by containing Ni. Therefore, when Cu and Ni are used in combination, the total amount is preferably 1. .5% or less, more preferably 1.0% or less.

[(D) B: more than 0% and 0.005% or less]
B is an element that improves hardenability, and is an element that is effective for allowing austenite to stably exist up to room temperature. In order to effectively exert such effects, the B content is preferably 0.0005% or more, more preferably 0.0010% or more, and still more preferably 0.0015% or more. However, if contained excessively, a boride is produced and ductility is deteriorated. Therefore, the B content is preferably 0.005% or less, more preferably 0.004% or less, and further preferably 0.0035% or less. To do.

[(E) Ca: at least one selected from the group consisting of more than 0% and 0.01% or less, Mg: more than 0% and 0.01% or less, and REM: more than 0% and 0.01% or less]
Ca, Mg, and REM are elements that have the effect of finely dispersing inclusions in the steel sheet, and may be contained alone or in combination of two or more selected arbitrarily. In order to effectively exhibit such an action, the contents of Ca, Mg, and REM are each preferably preferably 0.0005% or more, more preferably 0.0010% or more. However, when it is contained excessively, it causes deterioration of castability and hot workability. Accordingly, the contents of Ca, Mg, and REM are each preferably preferably 0.01% or less, more preferably 0.008% or less, and still more preferably 0.007% or less.

  In the present invention, REM is an abbreviation for rare earth elements, and includes lanthanoid elements, that is, 15 elements from La to Lu, and scandium and yttrium.

Next, a method for producing the high-strength cold-rolled steel sheet, high-strength electrogalvanized steel sheet, high-strength hot-dip galvanized steel sheet, and high-strength galvannealed steel sheet according to the present invention, particularly high-strength cold-rolled steel sheet and high-strength hot-dip zinc The manufacturing method of a plated steel plate is demonstrated. In addition, the manufacturing method of a high-strength cold-rolled steel sheet and the manufacturing method of a high-strength hot-dip galvanized steel sheet are “in a hot-rolling step of a steel sheet having the above-mentioned component composition, and wound at a winding temperature of 500 ° C. to 800 ° C. Hold at 500 ° C. or more and 800 ° C. or less for 3 hours or more, then cool to room temperature, and after cold rolling, keep soaked in a temperature range of (Ac ₁ point + 20 ° C.) to less than Ac ₃ point, and then average up to 500 ° C. The process until the cooling rate of 10 ° C./second or more and 500 ° C. or less is cooled to a temperature range of 500 ° C. or less at an average cooling rate of 10 ° C./second or more is the same. Since the reheating process after “cooling to a temperature range of 500 ° C. or lower” is different between the two, the process will be described separately for high-strength cold-rolled steel sheets and high-strength hot-dip galvanized steel sheets.

  In the manufacturing method of the high-strength cold-rolled steel sheet and the high-strength hot-dip galvanized steel sheet of the present invention, the steel sheet obtained by performing hot rolling and cold rolling on the steel satisfying the above component composition is subjected to annealing described later. Do. In the manufacturing method of a high-strength cold-rolled steel sheet, reheating is performed after the annealing. Furthermore, if necessary, a high-strength electrogalvanized steel sheet can be obtained by appropriately combining electrogalvanizing treatments. In the manufacturing method of a high-strength hot-dip galvanized steel sheet, after annealing, it is reheated and hot-dip galvanized. Furthermore, if necessary, a high-strength galvannealed steel sheet can be obtained by appropriately combining alloying treatments. In the present invention, a high-strength cold-rolled steel sheet or a high-strength hot-dip galvanized steel sheet having a desired structure can be obtained by appropriately controlling the manufacturing conditions.

For example, as shown in FIG. 1, hot rolling is performed based on a conventional method using steel having the above-described composition. In hot rolling, for example, after hot rolling so that the finish rolling temperature becomes Ac ₃ point or higher, winding is performed at a winding temperature of 500 ° C. or higher and 800 ° C. or lower. Thereafter, it is kept at 500 ° C. or more and 800 ° C. or less for 3 hours or more, then cooled to room temperature and cold-rolled. In addition, the cooling after finish rolling is about 500 ° C./second at the upper limit in operation.

After cold rolling, as an annealing process, soaking is maintained in a two-phase temperature range of Ac ₁ point + 20 ° C. or more and less than Ac ₃ point, and then cooled to 500 ° C. at an average cooling rate of 10 ° C./second or more. Cool at 500 ° C. or lower to a temperature range of 500 ° C. or lower at an average cooling rate of 10 ° C./second or higher. The method for producing a high-strength cold-rolled steel sheet includes a step of reheating to a temperature range of 250 ° C. or more and 500 ° C. or less, holding the temperature range for 30 seconds or more, and then cooling to room temperature. In the method for producing a high-strength hot-dip galvanized steel sheet, after cooling to the temperature range of 500 ° C. or lower, reheat to a temperature range of 250 ° C. or higher and 500 ° C. or lower and hold in the temperature range for 30 seconds or longer. A process of cooling to room temperature after hot dip galvanizing within the holding time is included.

  Hereinafter, the reason why each of the above conditions is specified will be described in detail.

[Winding at a winding temperature of 500 ° C. or higher and 800 ° C. or lower, and then holding at 500 ° C. or higher and 800 ° C. or lower for 3 hours or more, then cooling to room temperature]
After hot rolling, winding is performed at a winding temperature of 500 ° C. or higher and 800 ° C. or lower, and then held at 500 ° C. or higher and 800 ° C. or lower for 3 hours or longer to produce the predetermined Mn concentration distribution, and the carbide containing Mn It precipitates and becomes a hard phase spheroidized in the ferrite phase by annealing after cold rolling. In order to obtain such an effect, the coiling temperature is 500 ° C. or higher, preferably 550 ° C. or higher, more preferably 600 ° C. or higher. However, if the coiling temperature is too high, a large amount of scale or grain boundary oxidation occurs in the steel sheet and the pickling property deteriorates. Therefore, the coiling temperature is 800 ° C. or lower, preferably 750 ° C. or lower, more preferably 700 ° C. The following. The temperature range to be maintained after winding is 500 ° C. or higher, preferably 510 ° C. or higher, more preferably 520 ° C. or higher, still more preferably 550 ° C. or higher, and still more preferably 580 ° C. or higher. On the other hand, if the holding temperature is too high, the steel plate may cause a large amount of scale, grain boundary oxidation, etc., as in the case where the coiling temperature is too high, and the pickling property may deteriorate. Is 780 ° C. or lower, more preferably 750 ° C. or lower, and still more preferably 700 ° C. or lower. The holding time in the temperature range is 3 hours or more, preferably 4 hours or more, more preferably 5 hours or more, still more preferably 7 hours or more, and still more preferably 10 hours or more. On the other hand, if the holding time is too long, the holding time is preferably 72 hours because the steel sheet may cause a large amount of scale and grain boundary oxidation as well as the winding temperature is too high, and the pickling property may deteriorate. Below, more preferably 60 hours or less.

  In the present invention, holding at a predetermined temperature does not necessarily have to be held at the same temperature, and may vary as long as it is within a predetermined temperature range. For example, the temperature may be kept constant within the range of the holding temperature, and changes within this range, that is, temperature rise due to temperature drop or heating, temperature rise due to recuperation due to transformation, and the like are intended.

  In this invention, after hold | maintaining for the predetermined time in the said temperature range, it cools to room temperature, The cooling rate in that case is not specifically limited, For example, air cooling etc. may be sufficient.

[Pickling, cold rolling]
After hot rolling, pickling is performed as necessary, and cold rolling with a cold rolling rate of about 30 to 80% is performed.

[Annealing]
As an annealing process after cold rolling, soaking is maintained in a two-phase region of Ac ₁ point + 20 ° C. or more and less than Ac ₃ point, and then the temperature region up to 500 ° C. is cooled at an average cooling rate of 10 ° C./second or more. Then, the temperature range of 500 ° C. or lower is cooled at an average cooling rate of 10 ° C./second or higher, and the temperature range is cooled to 500 ° C. or lower.

The desired amount of ferrite can be ensured while maintaining the Mn concentration distribution of the present invention by maintaining soaking in a two-phase region of Ac ₁ point + 20 ° C. or higher and less than Ac ₃ point. If the soaking temperature is lower than Ac ₁ point + 20 ° C., the amount of ferrite in the metal structure of the finally obtained steel sheet increases so much that sufficient strength cannot be ensured. Therefore, the soaking temperature is Ac ₁ point + 20 ° C. or higher, preferably Ac ₁ point + 25 ° C. or higher, more preferably Ac ₁ point + 50 ° C. or higher, and further preferably Ac ₁ point + 80 ° C. or higher. On the other hand, when the Ac point is ₃ or more, ferrite cannot be sufficiently formed / grown during soaking, the ductility is lowered, the Mn concentration distribution is uniform, and the spherical hard phase formed in the ferrite phase is reduced. Decrease. For this reason, the soaking temperature is set to a temperature lower than Ac ₃ point, preferably Ac ₃ point −5 ° C. or lower, more preferably Ac ₃ point −10 ° C. or lower, and further preferably Ac ₃ point −20 ° C. or lower.

  In addition, the average temperature rising rate at the time of raising the temperature to the soaking temperature holding range is not particularly limited, and can be appropriately selected. For example, an average temperature rising rate of about 0.5 to 50 ° C./second may be used. .

  In the present invention, the holding time in the soaking temperature holding range is not particularly limited. However, if the holding time is too short, the processed structure remains and the ductility of the steel may be lowered. Therefore, the holding time is preferably 40 seconds or more, more preferably 60 seconds or more. On the other hand, if the retention time is too long, the concentration of Mn into the austenite phase proceeds, and the Ms point may decrease and the MA structure may increase. Therefore, the retention time is preferably 3600 seconds or less, more preferably 3000 seconds or less. To do.

In addition, as described above, in the present invention, holding at a predetermined temperature does not necessarily need to be held at the same temperature, and may be varied within a predetermined temperature range. For example, in the case of holding at the soaking temperature, the temperature may be kept constant within the range of Ac ₁ point + 20 ° C. or more and less than Ac ₃ point, or may be changed within this range.

The Ac ₁ point and Ac ₃ point can be calculated from the following formulas (a) and (b) described in “Leslie Steel Material Chemistry” (Maruzen Co., Ltd., issued May 31, 1985, page 273). . In the formula, [] indicates the content (% by mass) of each element, and the content of elements not included in the steel sheet may be calculated as 0% by mass.
Ac ₁ (° C.) = 723-10.7 × [Mn] −16.9 × [Ni] + 29.1 × [Si] + 16.9 × [Cr] + 290 × [As] + 6.38 × [W] · .. (a)
Ac ₃ (° C.) = 910−203 × √ [C] −15.2 × [Ni] + 44.7 × [Si] + 104 × [V] + 31.5 × [Mo] + 13.1 × [W] − ( 30 × [Mn] + 11 × [Cr] + 20 × [Cu] −700 × [P] −400 × [Al] −120 × [As] −400 × [Ti]) (b)

  After the soaking is maintained, the temperature range up to 500 ° C. is cooled at an average cooling rate of 10 ° C./second or more. By controlling the cooling rate from the soaking temperature, the generation of ferrite with a high Mn concentration can be suppressed and the amount of ferrite generated can be suppressed. When the average cooling rate is low, ferrite having a high Mn concentration is generated during cooling, which may deteriorate the bendability and the strength. Therefore, the average cooling rate is 10 ° C./second or more, preferably 15 ° C./second or more, more preferably 20 ° C./second or more. There is no particular upper limit on the average cooling rate, and water cooling or oil cooling may be used.

  After the temperature range up to 500 ° C. is cooled at the average cooling rate, 500 ° C. or lower is cooled at an average cooling rate of 10 ° C./second or higher. By controlling the average cooling rate of 500 ° C. or lower to 10 ° C./second or higher, the formation of soft high-temperature bainite can be suppressed, and the self-tempering of martensite can be suppressed to improve the strength. In order to obtain such an effect, the average cooling rate of 500 ° C. or lower is 10 ° C./second or higher, preferably 15 ° C./second or higher, more preferably 20 ° C./second or higher. There is no particular upper limit on the average cooling rate, and water cooling or oil cooling may be used.

  The average cooling rate up to 500 ° C. and the average cooling rate below 500 ° C. may be the same or different, and may be adjusted as appropriate within the above range.

  The cooling stop temperature when cooling the above 500 ° C. or lower at an average cooling rate of 10 ° C./second or higher is a temperature range of 500 ° C. or lower. When the cooling stop temperature is higher than 500 ° C., the hard phase is decreased, the strength cannot be ensured, and the MA structure increases and the bendability deteriorates. Therefore, the cooling stop temperature is 500 ° C. or lower, preferably 400 ° C. or lower, more preferably 350 ° C. or lower, and further preferably 300 ° C. or lower. The lower limit of the cooling stop temperature is not particularly limited, but is up to room temperature for operation.

  Hereinafter, the reheating process will be described separately for a method for manufacturing a cold-rolled steel sheet and a method for manufacturing a hot-dip galvanized steel sheet.

[Reheating in manufacturing method of cold-rolled steel sheet]
After stopping the cooling in the above temperature range of 500 ° C. or lower, reheating is performed to a temperature range of 250 ° C. or higher and 500 ° C. or lower, and the temperature is kept for 30 seconds or more and then cooled to room temperature.

  After the cooling is stopped, reheating is performed to a temperature range of 250 ° C. or more and 500 ° C. or less, and holding for 30 seconds or more can temper a hard phase such as martensite and transform untransformed austenite. When reheating is not performed or when the holding temperature is too low, tempering of the hard phase does not proceed, and high-density dislocations may be generated, or a large amount of MA structure may remain and bendability may be deteriorated. Therefore, the reheating temperature is 250 ° C. or higher, preferably 300 ° C. or higher, more preferably 350 ° C. or higher. On the other hand, if the holding temperature becomes too high, the strength decreases. Therefore, the reheating temperature is 500 ° C. or lower, preferably 470 ° C. or lower, more preferably 450 ° C. or lower. In the present invention, “reheating” means heating from the cooling stop temperature up to 500 ° C. or lower, that is, raising the temperature. Therefore, the reheating temperature is higher than the cooling stop temperature, and even in the temperature range of 250 ° C. or more and 500 ° C. or less, the isothermal holding or cooling stop where the cooling stop temperature and the reheating temperature are the same. The cooling process from temperature to lower temperature is not included in this reheating.

  If the holding time in the reheating temperature range is too short, the hard phase cannot be tempered sufficiently, and the untransformed austenite cannot be transformed. Accordingly, the holding time is 30 seconds or more, preferably 50 seconds or more, more preferably 100 seconds or more, and further preferably 200 seconds or more. On the other hand, the upper limit of the holding time is not particularly limited. However, if the holding time is too long, the productivity is lowered and the strength is lowered. Therefore, the holding time is preferably 1500 seconds or less, more preferably 1000 seconds or less.

  After holding for a predetermined time in the above reheating temperature range, it is cooled to room temperature. The average cooling rate at this time is not particularly limited, and is preferably 0.1 ° C./second or more, more preferably 0.4 ° C./second or more, preferably 200 ° C./second or less, more preferably 150 ° C./second or less. What is necessary is just to cool at an average cooling rate.

  In the present invention, an electrogalvanized layer (EG) may be formed on the obtained steel plate surface.

The formation method of said electrogalvanization layer is not specifically limited, The conventional electrogalvanization process method is employable. For example, in the case of producing an electrogalvanized steel sheet, a method of conducting an electrogalvanization process by energizing while being immersed in a zinc solution at 55 ° C. can be mentioned. Moreover, the plating adhesion amount per side is not particularly limited, and for example, in the case of an electrogalvanized steel sheet, it may be about 10 to 100 g / m ² .

[Reheating in hot-dip galvanized steel sheet manufacturing method]
After stopping cooling in the temperature range of 500 ° C. or lower, reheat to a temperature range of 250 ° C. or higher and 500 ° C. or lower, hold for 30 seconds or more, and apply hot dip galvanization within the holding time. Cool to room temperature.

  After the cooling is stopped, reheating is performed to a temperature range of 250 ° C. or more and 500 ° C. or less, and holding for 30 seconds or more can temper a hard phase such as martensite and transform untransformed austenite. When reheating is not performed or when the holding temperature is too low, tempering of the hard phase does not proceed, and high-density dislocations may be generated, or a large amount of MA structure may remain and bendability may be deteriorated. Therefore, the reheating temperature is 250 ° C. or higher, preferably 300 ° C. or higher, more preferably 350 ° C. or higher. On the other hand, if the holding temperature becomes too high, the strength decreases. Therefore, the reheating temperature is 500 ° C. or lower, preferably 470 ° C. or lower, more preferably 450 ° C. or lower. In the present invention, “reheating” means heating from the cooling stop temperature up to 500 ° C. or lower, that is, raising the temperature. Therefore, the reheating temperature is higher than the cooling stop temperature, and even in the temperature range of 250 ° C. or more and 500 ° C. or less, the isothermal holding or cooling stop where the cooling stop temperature and the reheating temperature are the same. The cooling process from temperature to lower temperature is not included in this reheating.

  If the holding time in the reheating temperature range is too short, the hard phase cannot be tempered sufficiently, and the untransformed austenite cannot be transformed. Accordingly, the holding time is 30 seconds or more, preferably 50 seconds or more, more preferably 100 seconds or more, and further preferably 200 seconds or more. On the other hand, the upper limit of the holding time is not particularly limited. However, if the holding time is too long, the productivity is lowered and the strength is lowered. Therefore, the holding time is preferably 1500 seconds or less, more preferably 1000 seconds or less.

  In the present invention, the hot dip galvanizing process is performed within a holding time of 30 seconds or more in the above reheating temperature region, and a hot dip galvanized layer is formed on the steel sheet surface. In the present invention, both hot dip galvanization and holding in the reheating temperature range are performed. That is, in order to appropriately manage the metal structure and strength by reheating, it is necessary to perform hot dip galvanization during the above holding time in the reheating temperature range. The method for forming the hot dip galvanized layer is not particularly limited, and a conventional hot dip galvanizing method can be employed. For example, the hot dip galvanizing process may be performed by immersing the steel sheet in a plating bath whose temperature is adjusted to the above reheating temperature range. The plating time only needs to satisfy the above holding time, and may be appropriately adjusted so as to ensure a desired plating amount. The plating time is preferably 1 to 10 seconds, for example.

There are the following various patterns as a combination of hot dip galvanizing treatment and reheating without plating treatment.
(I) After only heating, a hot dip galvanizing process is performed.
(Ii) After the hot dip galvanizing treatment, only heating is performed.
(Iii) Only heating, hot dip galvanization, and heating only are performed in this order.

  When the reheating temperature in the case of only heating is different from the hot dip galvanizing temperature, that is, the temperature of the plating bath, it may include the case of heating or cooling from one temperature to the other. Examples of the heating method include furnace heating and induction heating.

  When an alloyed hot dip galvanized layer is formed on the steel sheet surface, alloying may be performed after the hot dip galvanizing. Although the alloying temperature is not particularly limited, it is preferably 450 ° C. or higher, more preferably 460 ° C. or higher, and further preferably 480 ° C. or higher because alloying does not proceed sufficiently if the alloying temperature is too low. On the other hand, if the alloying temperature is too high, alloying proceeds too much and the Fe concentration in the plating layer increases and the plating adhesion deteriorates. Therefore, it is preferably 550 ° C. or less, more preferably 540 ° C. or less, and still more preferably. It is 530 degrees C or less. The time for the alloying treatment is not particularly limited, and may be adjusted so as to obtain a desired alloying. The alloying treatment time is preferably 10 seconds or more and 60 seconds or less. In addition, since alloying treatment is performed after holding for a predetermined time within the reheating temperature range, the alloying treatment time is not included in the holding time within the reheating temperature range.

  After holding for a predetermined time in the above reheating temperature range, it is cooled to room temperature. The average cooling rate at this time is not particularly limited, and is preferably 0.1 ° C./second or more, more preferably 0.4 ° C./second or more, preferably 200 ° C./second or less, more preferably 150 ° C./second or less. What is necessary is just to cool at an average cooling rate.

  The technique of the present invention can be suitably employed particularly for a thin steel plate having a thickness of 6 mm or less.

  The high-strength cold-rolled steel sheet and high-strength hot-dip galvanized steel sheet according to the present invention are intended for steel sheets having a tensile strength of 980 MPa or more, preferably 1,000 MPa or more, more preferably 1,010 MPa or more. The ductility is represented by a balance between strength and ductility (tensile strength (MPa) × ductility (%)), preferably 15,000 (MPa ·%) or more, more preferably 15,100 (MPa ·%) or more, Preferably, it is set to 15,200 (MPa ·%) or more. The bendability is expressed as a balance between strength and VDA bending angle (tensile strength (MPa) × VDA bending angle (°)), preferably 100,000 (MPa · °) or more, more preferably 100,000 (MPa · °). ) Or more, more preferably 101,000 (MPa · °) or more.

  EXAMPLES Hereinafter, the present invention will be described more specifically with reference to examples. However, the present invention is not limited by the following examples, but may be appropriately modified within a range that can meet the purpose described above and below. Of course, it is possible to implement them, and they are all included in the technical scope of the present invention.

[Example 1]
Steel having the composition shown in Table 1 below (the balance is iron and inevitable impurities, and the blank in Table 1 means that no element is added), and hot rolling → cold rolling → Continuous annealing was performed to produce a cold-rolled steel sheet.

[Hot rolling]
The slab was heated to 1250 ° C. and hot-rolled to a sheet thickness of 2.3 mm so that the reduction rate was 90% and the finish rolling temperature was 920 ° C. Then, after cooling from this temperature to the “winding temperature (° C.)” shown in Table 2 or Table 3 at an average cooling rate of 30 ° C./second, “holding temperature 1 (° C.)” shown in Table 2, And “holding time (hour)”, or held under the conditions of “holding start temperature (° C.)”, “holding end temperature (° C.)”, and “holding time (hour)” shown in Table 3. Subsequently, it air-cooled to room temperature and manufactured the hot-rolled steel plate.

[Cold rolling]
The obtained hot-rolled steel sheet was pickled to remove the scale on the surface, and then cold-rolled to produce a cold-rolled steel sheet having a thickness of 1.2 mm.

[Annealing of cold rolled steel sheet (CR)]
The obtained cold-rolled steel sheet was subjected to soaking maintenance → cooling → reheating under the conditions shown in Table 2 or Table 3 to produce a test steel. In Table 2, No. No. 32 is a comparative example in which reheating is not performed, and instead of reheating, after cooling from a cooling stop temperature of 480 ° C. to 350 ° C., the temperature is maintained for 300 seconds in the reheating column. .

  In the table, the soaking temperature is “soaking temperature (° C.)”, the average cooling rate up to 500 ° C. after soaking is “average cooling rate 1 (° C./sec)”, and the cooling rate at 500 ° C. or less is “ Average cooling rate 2 (° C / sec), cooling stop temperature is “cooling stop temperature (° C)”, holding temperature when reheating after cooling is “reheat holding temperature (° C)”, holding time at this holding temperature Is indicated as “reheat holding time (seconds)”. In this example, the holding time at the soaking temperature was set to 100 seconds to 600 seconds. After the “reheating holding temperature (° C.)” and the “reheating holding time (seconds)”, the steel was allowed to cool to room temperature to obtain a test steel. For cold-rolled steel sheets that were not electrogalvanized later, “CR” was entered in the “Product Type” column of the table.

[Manufacture of electrogalvanized steel sheet (EG)]
A part of the test steel was immersed in a galvanizing bath at 55 ° C., subjected to electroplating treatment (current density 30 to 50 A / dm ² ), washed with water and dried to obtain an electrogalvanized steel sheet. In addition, the amount of galvanized coating per side: 10 to 100 g / m ² . Further, in the above plating treatment, a test steel having an electrogalvanized layer on the surface was obtained by appropriately performing washing treatment such as alkaline aqueous solution degreasing, water washing and pickling. The electrogalvanized steel sheet was entered as “EG” in the “Type” column in the table.

  Each test steel was evaluated for metal structure, Mn concentration, and various mechanical properties as shown in detail below, and are shown in Table 4 or Table 5.

[Measurement of metal structure]
The area ratio of ferrite, the area ratio of the hard phase, the area ratio of the MA structure, the volume ratio of residual γ, and the ratio of the spherical hard phase were measured as follows. That is, after the cross section of the test steel was polished and corroded as shown below, a 1/4 position of the plate thickness was observed using an optical microscope or a scanning electron microscope (SEM). And the metal structure photograph image | photographed with the optical microscope or SEM was image-analyzed, and the ratio of each structure | tissue was measured. Details are shown below.

(Ferrite area ratio)
After the above polishing, it corrodes with nital, observes 3 fields of view (100 μm × 100 μm size / field of view) at a magnification of 1000 times with an SEM, and measures the area ratio of ferrite by a point calculation method with a lattice spacing of 5 μm and the number of lattice points of 20 × 20. The average value of 3 fields of view was calculated. The results are shown in “Ferrite (area%)” in the table. The area ratio of ferrite excludes the hard phase in the ferrite phase.

(Area ratio of hard phase)
The structure other than the ferrite was defined as a hard phase, and the value obtained by removing the ferrite area ratio from 100 area% of the observation visual field was defined as the area ratio of the hard phase. The results are shown in “Hard phase (area%)” in the table. The structure of the hard phase was also observed, and it was confirmed that the hard phase was at least one selected from the group consisting of bainitic ferrite, bainite, tempered martensite, residual γ, and MA structure.

(Area ratio of MA organization)
After the above polishing, it corrodes with a repeller, and is observed with an optical microscope at a magnification of 1000 at 3 fields of view (100 μm × 100 μm size / field of view), and the area ratio of the MA structure is determined by a point calculation method with a lattice spacing of 5 μm and a lattice number of 20 × 20. Measurements were made and the average value of 3 fields of view was calculated. The results are shown in “MA (area%)” in the table. In addition, the part whitened by the said repeller corrosion was observed as MA structure.

(Volume ratio of residual γ)
After polishing using # 1000 to # 1500 sandpaper to a thickness of 1/4, the surface is further electropolished to a depth of 10 to 20 μm, and then an X-ray diffractometer (RINT 1500 manufactured by Rigaku Corporation) is used. Measured. Specifically, a Co target is used, 40 kV-200 mA is output, a range of 40 ° to 130 ° is measured at 2θ, and the obtained diffraction peak (110), (200), (200) of ( 211) and fcc (γ) diffraction peaks (111), (200), (220), and (311), quantitative measurement of residual γ was performed. The results are shown in “Residual γ (volume%)” in the table.

(Spherical hard phase in ferrite phase)
After the polishing, the spherical hard phase corroded with nital and has an equivalent circle diameter of 2 μm or less and an aspect ratio of 1 to 3 present in the ferrite phase is observed with a SEM at a magnification of 1000 times (100 μm × 100 μm size / field of view). Then, the ratio of the spherical hard phase in the hard phase was determined by image analysis. The results are shown in “Spherical hard phase (area%)” in the table.

[Percentage of the region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel sheet]
The Mn concentration was measured by cutting the test steel in a cross section, embedding it in a resin, polishing, and polishing the range of 20 μm × 20 μm using EPMA under a beam diameter of 1 μm or less. The obtained Mn concentration was divided by the Mn concentration of the steel plate subjected to chemical analysis by inductively coupled plasma emission spectroscopy to determine the ratio of the region where the Mn concentration was concentrated 1.2 times or more with respect to the Mn concentration in the steel plate. . Thereafter, a region having a Mn concentration of 1.2 times or more and a region having a value of less than 1.2 times were color-coded, and the area% of the region having a Mn concentration of 1.2 times or more was determined. The results are shown in “Mn concentration 1.2 times area ratio (%)” in the table.

[Standard deviation of the region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel sheet]
The image color-coded according to the Mn concentration in the steel sheet is divided into 100 sections into 2 μm sections, and the fraction of the area where the Mn concentration is concentrated 1.2 times or more in each section is measured. Standard deviation was determined. The results are shown in “Standard deviation (area%) of Mn concentration 1.2 times region” in the table.

[Ratio of Mn concentration in ferrite phase to Mn concentration in steel sheet]
The same field of view as that of the above-mentioned Mn concentration measured by EPMA analysis of 20 μm × 20 μm was observed by SEM. Each ferrite grain and its Mn concentration distribution were identified by comparing the EPMA analysis result and the SEM image. The point where the major axis and the minor axis of the ferrite grain intersect was defined as the center position of the ferrite grain, and the Mn concentration at the center position was defined as the Mn concentration of the ferrite grain. By identifying the Mn concentration at the ferrite grain center position in the range of 20 μm × 20 μm by the above method and dividing the Mn concentration of the ferrite grain having the highest Mn concentration in each 20 μm × 20 μm range by the Mn concentration of the steel sheet, the ferrite phase The ratio of the Mn concentration in the Mn concentration of the steel sheet was determined. In the present invention, 20 μm × 20 μm is one field of view, and the average value of the three fields of view is the ratio of the Mn concentration in the ferrite phase to the Mn concentration of the steel sheet. The results are shown in the “Ratio of Mn concentration in ferrite phase” in the table.

[Evaluation of mechanical properties]
As for the mechanical properties of the test steel, a tensile test was performed using a No. 5 test piece defined by JIS Z2201, and the tensile strength and ductility were measured. The test piece was cut from the test steel so that the direction perpendicular to the rolling direction was the longitudinal direction. The strength-ductility balance was calculated from the obtained tensile strength and ductility. In the table, the tensile strength was “TS (MPa)”, the ductility was “EL (%)”, and the balance between strength and ductility was “TS × EL (MPa ·%)”.

  In the present invention, when the tensile strength was 980 MPa or more, the strength was high and the product was acceptable, and when the tensile strength was less than 980 MPa, the strength was insufficient and the product was evaluated as rejected.

  The ductility is evaluated by a strength-ductility balance. When TS × EL (MPa ·%) is 15,000 (MPa ·%) or more, the ductility is regarded as excellent, and when it is less than 15,000, the ductility is evaluated. It was evaluated as rejected as bad.

[Evaluation of bendability]
The bendability was evaluated under the following measurement conditions based on the VDA standard (VDA238-100) defined by the German Automobile Manufacturers Association. In the present invention, the displacement at the maximum load obtained by a bending test was converted into an angle based on the VDA, and the bending angle was obtained. The results are shown in “VDA bending angle (°)” in the table. The bendability was evaluated from the tensile strength and the bending angle. The results are shown in “TS × VDA (MPa · °)” in the table. When TS × VDA (MPa · °) was 100,000 (MPa · °) or more, it was determined that the bendability was excellent, and when it was less than 100,000, it was evaluated as unacceptable as insufficient bendability.
(Measurement condition)
Test method: roll support, punch push-in roll diameter: φ30mm
Punch shape: Tip R = 0.4mm
Distance between rolls: 2.9 mm
Pushing speed: 20mm / min
Specimen size: 60mm x 60mm
Bending direction: Rolling perpendicular direction Testing machine: SIMAZU AUTOGRAPH 20kN

  The following can be seen from Tables 1-5. Experiment No. manufactured with the annealing conditions prescribed | regulated by this invention using the steel types AT and W-AC which satisfy | fill the component composition of this invention. The steel sheets 1-18, 20, 24-26, 28, 29, and 33-43 were excellent in ductility and bendability in a region having a tensile strength of 980 MPa or more.

  On the other hand, as described in detail below, the steel sheets other than those described above did not satisfy the component composition and manufacturing conditions defined in the present invention, and the desired characteristics were not obtained.

  Steel type U in Table 1 had a C content, and steel type V had a Mn content exceeding the upper limit of the present invention, and fracture occurred during cold rolling, so the test steel could not be produced.

  Experiment No. No. 19 was an example in which the sample was cooled to 500 ° C. or lower and not reheated, and the MA structure increased and the bendability deteriorated.

  Experiment No. No. 21 was an example in which it was not kept at a predetermined temperature after winding, and the standard deviation of the Mn concentration 1.2 times region was low and the bendability was poor.

  Experiment No. No. 22 was an example in which the coiling temperature and the holding temperature were low. The area ratio with a Mn concentration of 1.2 times and the standard deviation of the region with a Mn concentration of 1.2 times were low, and the bendability was poor.

  Experiment No. No. 23 is an example in which the holding time at a predetermined temperature was short after winding, and the standard deviation of the Mn concentration 1.2 times region was low and the bendability was poor.

  Experiment No. No. 27 was an example in which the soaking temperature was high. Ferrite was not formed, and the area ratio with a Mn concentration of 1.2 times and the standard deviation of the region with a Mn concentration of 1.2 times were low, so the ductility was poor.

  Experiment No. No. 30 was an example in which the cooling stop temperature was high, and the ferrite and MA structures increased, the strength was low, and the ductility and bendability were also poor.

  Experiment No. No. 31 is an example in which the cooling rate to 500 ° C. is slow, and the bendability deteriorated because the Mn concentration in the ferrite phase became too high.

  Experiment No. No. 32 is an example in which reheating was not performed after cooling to 500 ° C. or lower, and the MA structure increased and the bendability deteriorated.

[Example 2]
Steel of the composition shown in Table 6 below (the balance is iron and inevitable impurities, and the blank in Table 6 means that no element is added), and hot rolling → cold rolling → Continuous annealing was performed to produce a cold-rolled steel sheet.

[Hot rolling]
The slab was heated to 1250 ° C. and hot-rolled to a sheet thickness of 2.3 mm so that the reduction rate was 90% and the finish rolling temperature was 920 ° C. Then, after cooling from this temperature to the “winding temperature (° C.)” shown in Table 7 or 8 at an average cooling rate of 30 ° C./second, “holding temperature 1 (° C.)” shown in Table 7, And “holding time (hour)”, or held under the conditions of “holding start temperature (° C.)”, “holding end temperature (° C.)” and “holding time (hour)” shown in Table 8. Subsequently, it air-cooled to room temperature and manufactured the hot-rolled steel plate.

[Cold rolling]
The obtained hot-rolled steel sheet was pickled to remove the scale on the surface, and then cold-rolled to produce a cold-rolled steel sheet having a thickness of 1.2 mm.

[Annealing of cold-rolled steel sheet (CR), manufacture of hot-dip galvanized steel sheet and alloyed hot-dip galvanized steel sheet]
The obtained cold-rolled steel sheet was subjected to soaking maintenance → cooling → reheating → plating treatment under the conditions shown in Table 9 or Table 10 to produce a test steel. In Table 7, No. 29 is a comparative example in which reheating is not performed. In the reheating column, instead of reheating, after cooling from a cooling stop temperature of 470 ° C. to 400 ° C., the temperature was maintained for 45 seconds. .

  In the table, the soaking temperature is “soaking temperature (° C.)”, the average cooling rate up to 500 ° C. after soaking is “average cooling rate 1 (° C./sec)”, and the cooling rate at 500 ° C. or less is “ Average cooling rate 2 (° C / sec), cooling stop temperature is “cooling stop temperature (° C)”, holding temperature when reheating after cooling is “reheat holding temperature (° C)”, holding time at this holding temperature “Reheat holding time (seconds)”, plating bath temperature “plating bath temperature (° C.)”, and plating treatment time “hot galvanizing treatment time (seconds)”, respectively. The “reheating holding time (seconds)” is the total time including the “hot dip galvanizing time (seconds)”.

  At the above-mentioned “reheating holding temperature (° C.)”, the steel sheet is immersed in a galvanizing bath after being held substantially (“reheating holding time (seconds)” — “hot dip galvanizing treatment time (seconds)”). The hot dip galvanized layer was formed with the “galvanizing treatment time (seconds)”. Experiment No. In Nos. 31 and 33, heating was performed from the “reheat holding temperature (° C.)” to “plating bath temperature (° C.)” in Table 8 immediately before dipping in the zinc plating bath, and then immersed in the zinc plating bath. Incidentally, some steel plates were subjected to galvanizing treatment and then alloying treatment. In the table, the alloying temperature at this time was expressed as “alloying temperature (° C.)”, and the holding time at the alloying temperature was expressed as “alloying time (seconds)”. After holding for a predetermined time, it was allowed to cool to room temperature to obtain a test steel. In the “product type” column of the table, “GI” was entered for the steel plate that was only subjected to the hot dip galvanizing treatment, and “GA” was entered for the steel plate that was also subjected to the alloying treatment.

  As described in detail below, each test steel was evaluated in the same manner as in Example 1 for the metal structure, Mn concentration, and various mechanical properties, and the results are shown in Table 9 or Table 10.

  The following can be seen from Tables 6-10. Experiment No. manufactured on the annealing conditions prescribed | regulated by this invention using the steel types AN and QU which satisfy | fill the component composition of this invention. The steel plates of 1, 2, 4, 5, 10, 11, 13-16, 18-23, 25-28, and 30-35 were excellent in ductility and bendability in a region having a tensile strength of 980 MPa or more.

  On the other hand, as described in detail below, the steel sheets other than those described above did not satisfy the component composition and manufacturing conditions defined in the present invention, and the desired characteristics were not obtained.

  Steel type O in Table 6 had a C content, and steel type P had a Mn content exceeding the upper limit of the present invention, and fracture occurred during cold rolling, so that the test steel could not be produced.

  Experiment No. No. 3 is an example in which the holding time at the reheating holding temperature after cooling to 500 ° C. or less was short. The MA structure increased and the bendability deteriorated.

  Experiment No. No. 6 was an example in which the coil was not held at a predetermined temperature after winding, and the standard deviation of the Mn concentration 1.2 times region was low and the bendability was poor.

  Experiment No. No. 7 was an example in which the coiling temperature and the holding temperature were low. The area ratio with a Mn concentration of 1.2 times and the standard deviation of the region with a Mn concentration of 1.2 times were low, and the bendability was poor.

  Experiment No. No. 8 was an example in which the holding time at a predetermined temperature was short after winding, the standard deviation of the Mn concentration 1.2 times region was low, and the bendability was poor.

  Experiment No. No. 9 is an example in which the reheating after cooling to 500 ° C. or lower was low. The MA structure increased and the bendability deteriorated.

  Experiment No. No. 12 was an example in which the soaking temperature was high, and ferrite was not sufficiently formed, and the Mn concentration in the ferrite phase was too high, so the ductility was poor.

  Experiment No. No. 17 was an example in which the cooling stop temperature was high, and the ferrite and MA structures increased, the strength was low, and the bendability was also poor.

  Experiment No. No. 24 is an example in which the cooling rate to 500 ° C. is slow, and the bendability deteriorated because the Mn concentration in the ferrite phase became too high.

  Experiment No. No. 29 was an example in which the sample was cooled to 500 ° C. or lower and not reheated, and the MA structure increased and the bendability deteriorated.

Claims (15)

The component composition of the steel sheet is mass%,
C: 0.10% or more and 0.30% or less,
Si: 1.2% or more and 3% or less,
Mn: 0.5% to 3.0%,
P: more than 0% and 0.1% or less,
S: more than 0% and 0.05% or less,
Al: 0.005% or more and 0.2% or less,
N: more than 0% and 0.01% or less, and O: more than 0% and 0.01% or less, the balance consisting of iron and inevitable impurities, and
A high-strength cold-rolled steel sheet having a tensile strength of 980 MPa or more excellent in ductility and bendability, wherein the structure of the steel sheet at a thickness of 1/4 position satisfies all of the following (1) to (5).
(1) When observed with a scanning electron microscope, the area ratio of ferrite to the entire structure is 5% or more and less than 50%, and the remainder is a hard phase.
(2) When the repeller corrosion is performed and observed with an optical microscope, the area ratio of the mixed structure of fresh martensite and retained austenite with respect to the entire structure is more than 0% and 30% or less.
(3) When analyzed with an electron beam microprobe analyzer, there are 5 area% or more of regions where the Mn concentration is 1.2 times or more the Mn concentration in the steel sheet, and (4) 2 μm sections The fraction of the region where the Mn concentration is concentrated 1.2 times or more of the Mn concentration in the steel sheet is measured, and the standard deviation when measuring 100 sections is 4.0% or more.
(5) When analyzed with an electron microprobe analyzer, the Mn concentration in the ferrite phase is 0.90 times or less of the Mn concentration in the steel sheet.   The high-strength cold-rolled steel sheet according to claim 1, wherein the volume ratio of retained austenite with respect to the entire structure is 5% or more when measured by an X-ray diffraction method.   The hard phase is composed of a mixed structure of the fresh martensite and retained austenite; and at least one structure selected from the group consisting of bainitic ferrite, bainite, and tempered martensite. High strength cold rolled steel sheet. The component composition is, as another element, in mass%,
The high-strength cold-rolled steel sheet according to any one of claims 1 to 3, comprising at least one selected from the group consisting of Cr: more than 0% and not more than 1% and Mo: more than 0% and not more than 1%. The component composition is, as another element, in mass%,
Ti: more than 0% and 0.15% or less,
The high-strength cold-rolled steel sheet according to any one of claims 1 to 4, containing at least one selected from the group consisting of Nb: more than 0% and 0.15% or less and V: more than 0% and 0.15% or less. . The component composition is, as another element, in mass%,
The high-strength cold-rolled steel sheet according to any one of claims 1 to 5, comprising at least one selected from the group consisting of Cu: more than 0% and 1% or less and Ni: more than 0% and 1% or less. The component composition is, as another element, in mass%,
The high-strength cold-rolled steel sheet according to any one of claims 1 to 6, containing B: more than 0% and 0.005% or less. The component composition is, as another element, in mass%,
Ca: more than 0% and 0.01% or less,
The high-strength cold-rolled steel sheet according to any one of claims 1 to 7, comprising at least one selected from the group consisting of Mg: more than 0% and not more than 0.01%, and REM: more than 0% and not more than 0.01%. .   A high-strength electrogalvanized steel sheet, wherein an electrogalvanized layer is formed on the surface of the high-strength cold-rolled steel sheet according to any one of claims 1 to 8.   A high-strength hot-dip galvanized steel sheet, wherein a hot-dip galvanized layer is formed on the surface of the high-strength cold-rolled steel sheet according to any one of claims 1 to 8.   A high-strength galvannealed steel sheet, characterized in that an alloyed galvanized layer is formed on the surface of the high-strength cold-rolled steel sheet according to any one of claims 1 to 8. A method for producing the high-strength cold-rolled steel sheet according to any one of claims 1 to 8,
In the hot rolling process of the steel sheet comprising the above component composition,
Winding is performed at a coiling temperature of 500 ° C. or more and 800 ° C. or less, then held at 500 ° C. or more and 800 ° C. or less for 3 hours or more, then cooled to room temperature, after cold rolling,
(Ac ₁ point + 20 ° C.) or higher and maintained at a temperature range of less than Ac ₃ point, then up to 500 ° C. with an average cooling rate of 10 ° C./second or more, 500 ° C. or less with an average cooling rate of 10 ° C./second or more, Cool to a temperature range below 500 ° C,
Next, a method for producing a high-strength cold-rolled steel sheet having a tensile strength of 980 MPa or more excellent in ductility and bendability, which is reheated to a temperature range of 250 ° C. or more and 500 ° C. or less, held for 30 seconds or more, and then cooled to room temperature.   A method for producing a high-strength electrogalvanized steel sheet, further comprising electrogalvanizing the high-strength cold-rolled steel sheet obtained by the production method according to claim 12. A method for producing the high-strength hot-dip galvanized steel sheet according to claim 10,
In the hot rolling process of the steel sheet comprising the above component composition,
Winding is performed at a coiling temperature of 500 ° C. or more and 800 ° C. or less, then held at 500 ° C. or more and 800 ° C. or less for 3 hours or more, then cooled to room temperature, after cold rolling,
(Ac ₁ point + 20 ° C.) or higher and maintained at a temperature range of less than Ac ₃ point, then up to 500 ° C. with an average cooling rate of 10 ° C./second or more, 500 ° C. or less with an average cooling rate of 10 ° C./second or more, Cool to a temperature range below 500 ° C,
Next, it is reheated to a temperature range of 250 ° C. or more and 500 ° C. or less, held for 30 seconds or more, and after being hot dip galvanized within the holding time, cooled to room temperature, has excellent tensile strength with excellent ductility and bendability. A method for producing a high-strength hot-dip galvanized steel sheet of 980 MPa or more.   The method for producing a high-strength hot-dip galvanized steel sheet according to claim 14, wherein alloying is performed in a temperature range of 450 ° C or higher and 550 ° C or lower after the hot-dip galvanizing.