high strength cold rolled steel sheet having excellent local deformability, and production method th

专利摘要:
Patent Summary: "High strength cold rolled steel plate having excellent local deformability, and method of production thereof". The present invention relates to this high strength cold rolled steel sheet containing,% by weight, c: 0.02% to 0.20%; si: 0.001% to 2.5%; mn: 0.01% to 4.0%; p: 0.001% to 0.15%; s: 0.0005% to 0.03%; al: 0.001% to 2.0%; n: 0.0005% to 0.01%; and o: 0.0005% to 0.01%; wherein si + al is limited to less than 1.0%, and the remainder being composed of iron and unavoidable impurities, wherein a bainite area ratio in a metal structure is 95% or more to a central portion. sheet thickness being a range of 5/8 to 3/8 in sheet thickness from the surface of the steel plate, an average pole density value of the orientation group {100} <011> to {223} < 110> is 4.0 or less, and a crystal orientation pole density {332} <113> is 5.0 or less, and an average crystal grain volume diameter in the metal frame is 7? M or any less.
公开号:BR112013026079B1
申请号:R112013026079-3
申请日:2012-04-12
公开日:2019-01-29
发明作者:Yoshihiro Suwa；Kazuaki Nakano；Kunio Hayashi；Riki Okamoto；Nobuhiro Fujita；Kohichi Sano
申请人:Nippon Steel & Sumitomo Metal Corporation；
IPC主号:

专利说明:

Report of the Invention Patent for "HIGH RESISTANCE COLD LAMINATED STEEL PLATE HAVING EXCELLENT LOCAL DEFORMABILITY, AND METHOD OF PRODUCTION".
Technical Field The present invention relates to a cold-rolled steel sheet which is excellent in local deformability, for bending, stretch bending, beveling operability, and the like, and is mainly used for automobile parts, and the like.
This application is based on, and claims the priority benefit of prior Japanese Patent Application No. 2011-089250, filed April 13, 2011, the contents of which are incorporated herein by reference. Background Art In order to suppress carbon dioxide emissions from vehicles, a reduction in the weight of a vehicle body has been promoted by the use of high strength steel plates. In addition, to also ensure passenger safety, a high strength steel plate is often used for the vehicle body in addition to a mild steel plate.
In order to further promote the reduction in the weight of motor vehicle bodies from now on, a wear resistance level of the high strength steel plate has to be increased more than conventionally, and in order to use the steel plate. High strength steel for a lower body part, for example, local bevel deformability has to be improved.
However, when a steel plate is increased in general strength, formability decreases, and, as shown in Non-Patent Document 1, uniform elongation important for stretching and bulging decreases. In contrast to this, as shown in Non-Patent Document 2, a method of ensuring uniform elongation even with the same strength by producing a metal structure of a complex steel plate is disclosed.
However, a method of controlling the metal structure of a steel plate that enhances the local ductility typified by bending, hole expansion, and chamfering is also disclosed. Non-Patent Document 3 reveals that control inclusions, which produce a uniform structure, and further decrease the hardness difference between the structures, are effective for bending and hole expansion.
This is to optimize the expansion of the bore by producing a uniform structure by controlling the structure, but in order to produce the uniform structure as shown in Non-Patent Document 4, a single phase austenite heat treatment becomes become a production base. Additionally, in order to achieve strength and ductility, Non-Patent Document 4 also discloses a technique in which metal structure control is performed by cooling control after hot rolling, precipitates are controlled, and a transformation structure is controlled. thereby obtaining appropriate fractions of ferrite and bainite.
However, Patent Document 1 discloses a method in which a finishing hot rolling temperature, a reduction ratio, and a finishing rolling temperature range are controlled, austenite recrystallization is promoted, development of a laminated texture. is suppressed, and crystal orientations are randomized, thereby improving hole resistance, ductility, and expansion. Prior Art Document Patent Document Patent Document 1: Japanese Published Patent Publication No. 2009-263718 Non-Patent Document Non-Patent Document 1: Kishida, Nippon Steel Technical Report (1999) No. 371, p. Non-Patent Document 2: O. Matsumura et al., Trans. ISIJ (1987) vol. 27, p. 570 Non-Patent Document 3: Kato et al., Steelmaking Research (1984) vol. 312, p. Non-Patent Document 4: K. Sugimoto et al., ISIJ International (2000) Vol. 40, p. 920 Description of the Invention Problems to be Solved by the Invention As described above, the main cause of local deformability deterioration is various "non-uniformities" of hardness difference between structures, non-metal inclusions, a developed laminate texture, and the like. The most effective non-uniformity between them is the hardness difference between the structures described in Non-Patent Document 3 above, and as another effective control factor, the developed laminate texture described in Patent Document 1 can be cited. These elements are mixed in a complex manner, and the local deformability of a steel sheet is determined. Therefore, to maximize an increased margin of local deformability by texture control, structure control is performed in a combined manner, and it is necessary to clearly eliminate non-uniformity for the hardness difference between structures.
Thus, the present invention is for providing a high strength cold rolled steel sheet having excellent local deformability capable of improving the local ductility of the high strength steel sheet, and also capable of perfecting anisotropy in the steel sheet by production. of a metal structure in which a bainite area ratio is 95% or more, together with a texture control, and a method of producing it.
Means to Solve Problems According to conventional knowledge, as described above, improvements in bore expansion, curvature, and the like have been accomplished by controlling inclusion, producing fine precipitates, homogenizing the structure, rotating structures in one phase. single, a decrease in hardness difference between the structures, and the like. However, these are not sufficient, so an effect on anisotropy is provided in a high strength steel plate to which Nb, Ti, and the like are added. This causes problems that other forming factors are sacrificed, the direction in which a blank before forming is taken is limited, and the like, and use is also limited.
Thus, the present inventors, in order to improve hole expansion and bending operability, have recently focused attention on the effect of a steel sheet texture, and have examined and studied its functional effect in detail. As a result, they revealed that by controlling the intensities of respective orientations of a specific crystal orientation group, local deformability improves dramatically without elongation and strength decreasing greatly. The point where emphasis should be placed is that they also revealed that an improved margin of local deformability by texture control occurs greatly in a steel structure, a metal structure in which a bainite area ratio of 95% or more is produced. and thus the improved margin of local deformability is maximized on the basis that steel strength is assured. Additionally, they found that in a structure in which the respective orientation intensities of a specific crystal orientation group are controlled, the size of a grain unit greatly affects local ductility.
Generally, in a structure in which low temperature generation phases (bainite, martensite, and the like) are mixed, the definition of crystal grains is extremely vague, and quantifying them is difficult. In contrast, the present inventors have found it possible to solve the problem of crystal grain quantification if a "grain unit" of crystal grain is determined in the following manner. The crystal grain "grain unit" determined in the present invention is determined as follows in an orientation analysis of a steel plate by an Electron Back Scattering Pattern (EBSP). That is, in an orientation analysis of a sheet steel by an EBSP, for example, orientations are measured at 1500 magnifications with a measured step of 0.5 μηη or less, and a position in which a misorientation between adjacent measured points exceeding 15 ° is set to a boundary between crystal grains. Then a region surrounded with this limit is determined to be the "grain unit" of crystal grains.
With respect to the crystal grains of the grain unit determined in this way, a circle equivalent diameter d is obtained, and the crystal grain volume of each grain unit is obtained by 4 / 3Kd3. Then a weighted average volume is calculated, and an average volume diameter (Average Volume Diameter) is obtained. The present invention is based on the knowledge previously described, and the essence of it is as follows.
[1] A high-strength cold-rolled steel plate having excellent local deformability contains: by weight%, C: not less than 0,02% and not more than 0,20%;
Si: not less than 0,001% or more than 2,5%;
Mn: not less than 0,01% and not more than 4,0%; P: Not less than 0.001% and not more than 0.15%; S: not less than 0.0005% nor more than 0.03%;
Al: not less than 0.001% nor more than 2.0%; N: not less than 0.0005% nor more than 0.01%; and O: not less than 0.0005% nor more than 0.01%; wherein Si + Al is limited to less than 1.0%, and the remainder being composed of iron and unavoidable impurities, where a bainite area ratio in a metal structure is 95% or more in a central portion. sheet thickness being a range of 5/8 to 3/8 in sheet thickness from the surface of the steel plate, an average pole density value of the orientation group {100} <011> to {223} < 110> represented by their crystal orientations of {100} , {116} <110>, {114} <110>, {113} <110>, {112} <110>, {335} <110> , and {223} <110> is 4.0 or less, and a crystal orientation pole density {332} <113> is 5.0 or less, and an average crystal grain volume diameter in the structure of metal is 7 μηη or less.
[2] High strength cold-rolled steel plate having excellent local deformability according to [1], where bainite crystal grains, a ratio of crystal grains where a ratio of a length dL in A rolling direction for a length dt in a plate thickness direction: dL / dt is 3.0 or less is 50% or more.
[3] High-strength cold-rolled steel plate having excellent local deformability according to [1] additionally contains: one or two or more types of mass%, Ti: not less than 0.001% nor more than 0.20%, Nb: not less than 0.001% or more than 0.20%, V: not less than 0.001% or more than 1.0%, and W: not less than 0.001 % not more than 1,0%.
[4] High-strength cold-rolled steel plate having excellent local deformability according to [1] additionally contains: one or two or more types of bulk%, B: not less than 0.0001 % not more than 0,0050%, Mo: not less than 0,001% or more than 1,0%, Cr: not less than 0,001% or more than 2,0%, Cu: not less than 0.001% not more than 2.0%, Ni: not less than 0.001% or more than 2.0%, Co: not less than 0.0001% or more than 1.0%, Sn: no less than 0.0001% not more than 0.2%, Zr: not less than 0.0001% not more than 0.2%, and As: not less than 0.0001% not more than 0.50%.
[5] High strength cold rolled steel sheet having excellent local deformability according to [1] additionally contains: one or two or more types of bulk%, Mg: not less than 0.0001 % not more than 0.010%, REM: not less than 0.0001% not more than 0.1%, and Ca: not less than 0.0001% not more than 0.010%.
[6] High-strength cold-rolled steel sheet having excellent local deformability according to [1], wherein on the surface a hot-dip galvanized layer or an alloy hot-dip galvanized layer is provided.
[7] A method of producing a high-strength cold-rolled steel plate having excellent local deformability includes: in a steel billet containing: by weight%, C: not less than 0,02% or more than 0.20%;
Si: not less than 0,001% or more than 2,5%;
Mn: not less than 0,01% and not more than 4,0%; P: Not less than 0.001% and not more than 0.15%; S: not less than 0.0005% nor more than 0.03%;
Al: not less than 0.001% nor more than 2.0%; N: not less than 0.0005% nor more than 0.01%; and O: not less than 0.0005% nor more than 0.01%; where Si + Al is limited to less than 1.0%, and the remainder being composed of iron and unavoidable impurities, first hot rolling wherein lamination at a reduction ratio of 40% or more is performed once or more in a temperature range of not lower than 1000 ° C or higher than 1200 ° C; adjusting an austenite grain diameter to 200 μηη or less by the first hot rolling; second hot rolling wherein lamination at a reduction ratio of 30% or more is performed in one pass at least once in a temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C, as determined by Expression (1) below; total reduction ratio adjustment on second hot rolling to 50% or more; performing final reduction at a reduction ratio of 30% or more on the second hot rolling and then initiating primary cooling such that a wait time t second satisfies Expression (2) below; adjusting an average primary cooling rate at 50 ° C / second or more, and performing primary cooling such that a temperature change is within a range of not lower than 40 ° C, or higher than that 140 ° C; performing cold rolling at a reduction ratio of not less than 30%, not more than 70%; holding for 1 to 300 seconds / seconds in a temperature region of Ae3 to 950 ° C; performing secondary cooling at an average cooling rate of not less than 10 ° C / s or more than 200 ° C / s in a temperature region of Ae3 to 500 ° C; and conducting an excess heat treatment wherein retention is performed for not less than t2 seconds satisfying Expression (4) below, not more than 400 seconds in a temperature region of not lower than 350 ° C or higher. high than 500 ° C. T1 (° C) = 850 + 10 x (C + N) x Mn + 350 x Nb + 250 x Ti + 40 x B + 10xCr + 100xMo + 100 xV ··· (1) t ^ 2.5 xt1 ··· (2) Here, t1 is obtained by Expression (3) below. t1 = 0.001 x ((Tf - T1) x P1 / 100) 2 - 0.109 x ((Tf - T1) x P1 / 100) + 3.1 ··· (3) Here, in Expression (3) above, Tf represents the temperature of the steel billet obtained after the final reduction by a reduction ratio of 30% or more, and P1 represents the reduction ratio of the final reduction to 30% or more.
Iog (t2) = 0.0002 (T2 - 425) 2 + 1.18 ... (4) Here, T2 represents an excess treatment temperature, and the maximum value of t2 is set to 400.
[8] The method of production of high strength cold rolled steel sheet having excellent local deformability according to [7], wherein the total reduction ratio in a temperature range of lower than T1 + 30 ° C is 30% or less.
[9] The method of producing the high strength cold rolled steel plate having excellent local deformability according to [7], wherein the wait time t second additionally satisfies Expression (2a) below. t <t1 ··· (2a) [10] The method of producing high-strength cold-rolled steel plate having excellent local deformability according to [7], where the holding time t second additionally satisfies the Expression (2b) below. t1 ^ t ^ t1x2,5- (2b) [11] The method of producing high-strength cold-rolled steel plate having excellent local deformability according to [7], wherein primary cooling is initiated between supports. Lamination
[12] The method of producing high strength cold rolled steel sheet having excellent local deformability according to [7], where when heating is carried out to the temperature region of Ae3 at 950 ° C after rolling. in cold weather, an average heating rate of not lower than room temperature or higher than 650 ° C is adjusted to HR1 (° C / second) expressed by Expression (5) below, and an average heating rate from higher than 650 ° C to Ae3 at 950 ° C is adjusted to HR2 (° C / second) expressed by Expression (6) below. HR1 ^ 0.3 ... (5) HR2 ^ 0.5x HR1 ... (6) [13] The production method of high strength cold rolled steel sheet having excellent local deformability according to [ 7] further includes: forming a hot dip galvanized layer or a hot dip galvanized alloy layer on the surface.
Effect of the Invention In accordance with the present invention, it is possible to obtain a high strength cold-rolled steel sheet having excellent local deformability for bending, drawing flanging, chamfering, and the like because it contains a steel sheet texture and structure. of steel.
Brief Description of the Drawings FIG. 1 shows the relationship between an average pole density value of the orientation group {100} To {223} <110> and a minimum plate thickness / bending radius; FIG. 2 shows the relationship between a crystal orientation pole density {332} <113> and the plate thickness / minimum bending radius; FIG. 3 shows the relationship between the number of times of rolling at 40% or more in rough rolling and austenite grain diameter in rough rolling; FIG. 4 shows the relationship between the reduction ratio at T1 + 30 at T1 + 200 ° C and the average pole density value of the orientation group {100} <011> a {223} <110>; FIG. 5 shows the relationship between the reduction ratio at T1 + 30 to T1 + 200 ° C and the crystal orientation pole density {332} <113>; FIG. 6 is an explanatory view of a continuous hot rolling line; FIG. 7 shows the relationship between strength and bore expansion of the steels of the present invention and comparative steels; and FIG. 8 shows the relationship between the strength and curvature of the steels of the present invention and the comparative steels.
Mode of Carrying Out the Invention Hereinafter, the contents of the present invention will be explained in detail. (Crystal Orientation) First, an average pole density value of the {100} To {223} <110> orientation group and a crystal orientation pole density {332} <113> to a central portion of sheet thickness being a range of 5/8 to 3/8 in sheet thickness of a surface of a steel sheet.
In a co-rolled sheet steel of the present invention, an average pole density value of the {100} To {223} <110> orientation group and a pole density of the {332} <113 orientation group > a central plate thickness portion being a range of 5/8 to 3/8 plate thickness from the surface of the steel plate are particularly important characteristic values.
As shown in FIG. 1, when X-ray diffraction is performed on the central plate thickness portion with the range of 5/8 to 3/8 in plate thickness from the steel plate surface to obtain pole densities of respective orientations, the mean pole density value of the {100} to {223} <110> orientation group is less than 4.0, and it is possible to satisfy a sheet thickness / bending radius 1.5 which is required to operate a piece of structure to be recently required. Additionally, when in a steel structure, 95% or more of a bainite fraction is satisfied, the sheet thickness / bending radius 2.5 is satisfied. When hole expansion is small, limited curvature is required, the average value of the pole densities of the {100} To {223} <110> orientation group is desirably less than 3.0.
When this value is 4.0 or higher, the anisotropy of the mechanical properties of the steel plate becomes extremely strong, and additionally local deformability in only one direction is improved, but a material in a different direction of the same deteriorates significantly. it, resulting in it becoming impossible to satisfy the plate thickness / bending radius 1.5. On the other hand, when this value becomes less than 0.5, which is difficult to achieve in a current general continuous hot rolling process, deterioration of local deformability is provided.
The guidelines {100} <011>, {116} <110>, {114} <110>, {113} <110>, {112} <110>, {335} <110>, and {223} <110 > are included in this guidance group. Pole density is synonymous with a random X-ray intensity ratio. Pole density (random x-ray intensity ratio) is a numerical value obtained by measuring x-ray intensities of a standard sample having no accumulation in a specific orientation, and a test sample under the same conditions by diffractometry. x-rays or similar, and dividing the x-ray intensity of the test sample by the x-ray intensity of the standard sample. This pole density can be measured by either X-ray diffraction, an Electron Back Scattering Pattern (EBSP) method, and an Electron Channeling Pattern (ECP) method.
As for the pole density of the orientation group {100} <011> to {223} <110>, for example, pole densities of respective orientations of {100} <011>, {116} <110>, {114 } <110>, {112} <110>, and {223} <110> are obtained from a three-dimensional texture (ODF) calculated by a series expansion method using a plurality (preferably three or more) of outside pole figures. from the pole figures of {110}, {100}, {211}, and {310} measured by the method, and these pole densities are arithmetic means, and thus the pole density of the orientation group described above is obtained. Incidentally, when it is impossible to obtain the intensities of all orientations described above, the arithmetic mean of the pole densities of the respective orientations of {100}, {116} <110>, {114} <110>, {112} <110>, and {223} <110> can also be used as a substitute.
For example, for the pole densities of each of the crystal orientations described above, each of the intensities of (001) [1-10], (116) [1-10], (114) [1-10], (113) [1-10], (112) [1-10], (335) [1-10], and (223) [1-10] at a cross section of φ2 = 45 ° in the three-dimensional texture can be used as it is.
Additionally, due to the similar reason, the pole density of the orientation group {332} <113> of the sheet plane in the central portion of sheet thickness being a range of 5/8 to 3/8 in sheet thickness from the The surface of the steel sheet must be 5.0 or less as shown in FIG. 2. Considering that it is desirably 3.0 or less, the sheet thickness / bending radius 1.5 that is required to operate a newly required frame portion is satisfied. Additionally, when in the steel structure, 95% or more of the bainite fraction is satisfied, the sheet thickness / bending radius ^ 2.5 is satisfied. On the other hand, when the pole density of the {332} <113> orientation group is greater than 5.0, the anisotropy of the steel sheet mechanical properties becomes extremely strong, and in addition the local deformability only In a certain direction it is improved, but material in a different direction deteriorates significantly, resulting in it being impossible to safely satisfy the sheet thickness / bending radius ^ 1,5. Additionally, when the pole density becomes less than 0.5, which is difficult to achieve in a current general continuous hot rolling process, deterioration of local deformability is provided. The reason why the pole densities of the crystal orientations described above are important for freezing property of the shape at the bending operation is not necessarily obvious, but is inferentially related to the irregular behavior of the crystal at the bending deformation.
With respect to the sample being subjected to X-ray diffraction, EBSP method, or ECP method, the steel plate is reduced in thickness to a predetermined plate thickness from the surface by mechanical polishing, or the like. . Thereafter, the resistance is removed by chemical polishing, electrolytic polishing, or the like, and the sample is manufactured in such a way that in the range of 5/8 to 3/8 plate thickness, an appropriate plane becomes a measurement. For example, on a steel part in a size of 30 mm φ cut from the 1/4 W or 3/4 W width of the W plate, fine-finish milling (average centerline roughness Ra: 0 , 4a to 1.6a) is performed. Then, by chemical polishing or electrolytic polishing, the resistance is removed, and the sample to be subjected to X-ray diffraction is fabricated. With respect to the width direction of the plate, the steel part is desirably taken from the steel plate, the 1/4 or 3/4 position of an end portion.
Of course, the pole density satisfying the limited range of pole density described above not only in the central plate thickness portion being a range of 5/8 to 3/8 in plate thickness from the steel plate surface, but also in many possible thickness positions, and thus the performance of local ductility (local elongation) is further improved. However, the range of 5/8 to 3/8 from the surface of the steel sheet is measured, thereby making it possible to represent the property of the overall steel sheet material generally. Thus, 5/8 to 3/8 of the plate thickness is prescribed as the measuring range.
Incidentally, the crystal orientation represented by {h-kl} <uvw> means that the normal direction of the sheet steel plane is parallel to <hkl>, and the rolling direction is parallel to <uvw>. With respect to the crystal orientation, typically the vertical plane plane orientation is represented by [hkl] or {hkl}, and the orientation parallel to the rolling direction is represented by (uvw) or <uvw>. {hkl} and <uvw> are generic terms for equivalent planes, and [hkl] and (uvw) each indicate an individual crystal plane. That is, in the present invention, a body-centered cubic structure is objectified, and thus, for example, the planes (111), (-111), (1-11), (11-1), (-1 -11), (-11-1), (1-1-1), and (-1-1-1) are equivalent to make it impossible to make them different. In such a case, these guidelines are generally referred to as {111}. In an ODF representation, [hkl] (uvw) is also used to represent orientations of other low symmetrical crystal structures, and thus is general to represent each orientation as [hkl] (uvw), but in the present invention. , [hkl] (uvw) and {hkl} <uvw> are synonymous with each other. X-ray crystal orientation measurement is performed according to the method described in, for example, Cullity, Elements of X-ray Diffraction, new addition (published in 1986, translated by MATSUMURA, Gentaro, published by AGNE Inc. ) on pages 274 to 296. (Average Crystal Grain Volume Diameter) The present inventors have diligently examined the texture control of a hot-rolled steel sheet. As a result, it has been found that under the condition that a texture is controlled as described above, the effect of crystal grains in one grain unit on local ductility is extremely large, and crystal grains are produced thin, thereby rendering drastic improvement of local ductility can be achieved. Incidentally, as described above, the "grain unit" of the crystal grains is determined such that the position at which a misorientation exceeds 15 ° is adjusted as a crystal grain limit in a steel sheet orientation analysis by EBSP.
As above, the reason why local ductility perfects is not obvious. However, it is conceivable because when the texture of the steel plate is randomized and the crystal grains are produced thin, the local strength concentration to occur in the micron order is suppressed, the strain homogenization is increased, and the strength is uniformly dispersed. in micron order.
As there are no longer large crystal grains even though their number is small, the deterioration of local ductility becomes greater. Therefore, crystal grain size is not an ordinary average size, and an average volume diameter defined as a weighted average volume is strongly correlated with local ductility. In order to achieve this effect, the average volume diameter of the crystal grains needs to be 7 μηη or less. It is desirably 5 μηη or less to ensure that the hole expands at a higher level. Incidentally, the crystal grain measurement method is adjusted as described above. (Equiaxial property of crystal grains) As a result of further seeking local ductility, the present inventors have also found that when the equiaxial property of crystal grains is excellent on condition that the texture described above and the size of the crystal grains are satisfied, local ductility perfects. As an index indicating this equiaxial property, in relation to crystal grains expressed per grain unit, an excellent grain ratio in equiaxial property where dL / dt, being a crystal grain ratio, a length dL in a rolling direction. cold at a length dt in one direction of plate thickness, is 3.0 or less, needs to be at least 50% or more for all bainite grains. (Chemical Composition) Subsequently, the limiting conditions of the components will be described. Incidentally, the% of each content is mass%. C: Not less than 0.02%, not more than 0.20% The lower limit of C is adjusted to 0.02% to have 95% or more of bainite in the steel structure. Additionally, C is a strength enhancing element, thereby preferably being adjusted to 0.025% or more to ensure strength. On the other hand, when the C content exceeds 0.20%, weldability is sometimes impaired, and operability sometimes deteriorates greatly due to an increase in a rigid structure, and thus the upper limit is set to 0. , 20%. Additionally, when the C content exceeds 0.10%, the formability deteriorates, so that the C content is preferably adjusted to 0.10% or less.
Si: Not less than 0.001%, not more than 2.5% Si is an effective element for increasing the mechanical strength of the steel sheet, but when Si becomes greater than 2.5%, the operability deteriorates, and a notch of the surface occurs, so that it is set to the upper limit. Additionally, when the Si content is large, a chemical conversion treatment property decreases, so that Si is preferably adjusted to 1.20% or less. On the other hand, it is difficult to adjust Si to less than 0.001% in a practical steel, so that it is set to the lower limit.
Mn: Not less than 0.01%, not more than 4.0% Mn is also an effective element for increasing the mechanical strength of the steel sheet, but when Mn becomes greater than 4.0%, operability deteriorates so that it is adjusted to the upper limit. On the other hand, it is difficult to adjust Mn to less than 0.01% on a practical steel, so that it is adjusted to the lower limit. Additionally, when elements such as Ti which suppresses the occurrence of hot cracking caused by S are not sufficiently added except Mn, the amount of Mn satisfying Mn / S 20% is desirably added. Additionally, Mn is an element that, with an increase in grade, expands an austenite region temperature to a low temperature side, improves hardness, and facilitates the formation of a continuous cooling transformation structure having excellent beveling operability. This effect is not easily displayed when the Mn content is less than 1%, so that 1% or more is desirably added. P: not less than 0.001%, not more than 0.15% S: not less than 0.0005%, not more than 0.03% With respect to the upper limits of P and S, P is adjusted to 0.15% or less, and S is adjusted to 0.03% or less, respectively. This is to prevent deterioration of operability and cracking at the time of hot rolling or cold rolling. With respect to the lower limits of P and S, P is set to 0.001%, and S is set to 0.0005% as applicable values for current general refining (including secondary refining).
Al: Not less than 0.001%, not more than 2.0% For deoxidation, 0.001% or more of Al is added. When deoxidation is sufficiently required, 0.01% or more is preferably added. Additionally, Al is also a significant element that increases a transformation point ya a. However, when it is much larger, weldability deteriorates so that the upper limit is set at 2.0%. It is preferably adjusted to 1.0% or less. N: not less than 0.0005%, not more than 0.01% O: not less than 0.0005%, not more than 0.01% N and O are impurities, and are both adjusted to 0 , 01% or less to prevent deterioration operability. The lower limits of both elements are set to 0.0005% which is applicable to current general refining (including secondary refining). However, they are preferably adjusted to 0.001% or more to suppress an extreme increase in the cost of steel production.
Si + Al: less than 1.0% When Si and Al are excessively contained, precipitation of cementite during an over treatment is suppressed and the retained austenite fraction becomes very large, so that the total amount added of Si and Al is adjusted to less than 1%.
Ti: not less than 0.001% or more than 0.20% Nb: not less than 0.001% or more than 0.20% V: not less than 0.001% or more than 1.0% W: not less than 0.001% nor more than 1.0% Additionally, when resistance is obtained by precipitation resistance, it is preferred to generate fine carbonitride. To obtain precipitation resistance, it is effective to add Ti, Nb, V, and W, and one or two or more types thereof may be contained.
In order to achieve this effect by adding Ti, Nb, V, and W, it is necessary to add 0.001% Ti, 0.001% Nb, 0.001% or more of V, and 0.001% or more of W. When resistance to precipitation is particularly necessary, it is desired to add 0.01% or more of Ti, 0.005% or more of Nb, 0.01% or more of V, and 0.01% or more of W. However, even when they are added excessively, the resistance increase is saturated, and additionally recrystallization after hot rolling is suppressed, thus making it difficult to perform crystal orientation control after cold rolling annealing so that Ti needs to be set to 0 , 20% or less, Nb needs to be adjusted to 0.20% or less, V needs to be adjusted to 1.0% or less, and W needs to be adjusted to 1.0% or less. B: not less than 0.0001%, not more than 0.0050% Mo: not less than 0.001%, not more than 1.0% Cr: not less than 0.001%, not more than 2 , 0% Cu: not less than 0.001%, not more than 2.0% Ni: not less than 0.001%, not more than 2.0% Co: not less than 0.0001%, not more 1.0% Sn: not less than 0.0001%, not more than 0.2% Zr: not less than 0.0001%, not more than 0.2% As: no less than 0.0001%, not more than 0.50% When strength is ensured by increasing the hardness of the structure to perform the second phase control, it is effective to add one or two or more types of B, Mo, Cr, Cu. , Ni, Co, Sn, Zr, and As. In order to achieve this effect, it is necessary to add 0.0001% or more of B, 0.001% or more of each Mo, Cr, Cu, and Ni, and 0, 0001% or more of each of Co, Sn, Zr, and As. However, when they are added excessively, operability is deteriorated otherwise, so that the upper limit of B is set to 0.0050%, the upper limit of Mo is set to 1.00%, upper limit of each of Cr, Cu, and Ni is adjusted to 2.0%, upper limit of Co is adjusted to 1.0%, upper limit of each Sn and Zr is set to 0.2%, and the upper limit of As is set to 0.50%.
Mg: not less than 0.0001%, not more than 0.010% REM: not less than 0.0001%, not more than 0.1% Ca: not less than 0.0001%, not more than 0.010% Mg, REM, and Ca are important elements to be added to improve local formability and produce harmful inclusions. In order to achieve this effect, the lower limit of each is set to 0.0001%. On the other hand, excessive additions lead to cleaning deterioration, so that 0.010% is set as the upper limit of Mg, 0.1% is set as the upper limit of REM, and 0.010% is set as the upper limit of Ca (Metal Structure) In the following, a cold-rolled steel sheet metal structure of the present invention will be explained here. The metal structure of the cold rolled steel sheet of the present invention has a bainite area ratio of 95% or more, and is preferably a single bainite phase. This is because the metal structure is composed of bainite, thereby making it possible to achieve hole strength and expansion. Additionally, this structure is generated by transformation at a relatively high temperature, so as not to have to be cooled at low temperature when being produced, and is a preferred structure also in terms of material stability and productivity.
Like the rest, 5% or less of proeutectoid ferrite, pearlite, martensite, and retained austenite is allowed. Proeutectoid ferrite is not a problem, considering that it is sufficiently resistant to precipitation, as proeutectic ferrite sometimes becomes soft depending on the chemical composition, and when the area ratio becomes larger than than 5%, the hole expansion slightly decreases due to the hardness difference of the bainite. Additionally, when a pearlite area ratio becomes greater than 5%, strength and operability sometimes deteriorate. When the ratios of martensite and retained austenite area to be transformed induced by resistance to martensite become 1% or more, and greater than 5%, respectively, an interface between bainite and a structure hardener than bainite becomes. It becomes a cracking starting point, and the expansion of the hole deteriorates.
Thus, considering that the area ratio of bainite is set at 95% or more, the area ratio of proeuctoid ferrite, pearlite, martensite, and γ retained being the remainder becomes 5% or less, so the strength and expansion of the hole are well supported. However, it is necessary to adjust the martensite to less than 1% as described above.
Here, the bainite in the present invention is a microstructure defined as a continuous cooling transformation structure (Zw) positioned at an intermediate stage between a microstructure containing polygonal ferrite and pearlite to be generated by a diffusion mechanism, and martensite to be generated by a non-diffusive shear mechanism as described in The Iron and Steel Institute of Japan, Society of Basic Research, Bainite Research Committee / Edition; Recent Research on Bainitic Microstructures and Transformation Behavior of Low Carbon Steels - Final Report of Bainite Research Committee (in 1994, The Iron and Steel Institute of Japan).
That is, the continuous cooling transformation structure (Zw) is defined as a microstructure mainly composed of bainitic ferrite (oc ° b), granular bainitic ferrite (aB), and quasi-polygonal ferrite (otq), and additionally containing a small amount of retained austenite (γΓ) and Martensite-austenite (MA) as described in the reference literature described above on pages 125 to 127 as an optical microscope observation structure.
Incidentally, similar to polygonal ferrite (PF), an internal structure of otq does not appear by engraving, but a form of otq is acicular, and is definitely distinguished from PF. Here, on the condition that of an object crystal grain, a peripheral length is set to Iq and an equivalent circle diameter is set to dq, a grain having a ratio (Iq / dq) of them satisfying Iq / dq ^ 3.5 is otq. The continuous cooling transformation structure (Zw) of the present invention is defined as a microstructure containing one or two or more types of ot ° B, otB, otq, γΓ, and MA. Incidentally, the total content of γΓ and MA being small in quantity is adjusted to 3% or less.
Here is sometimes a case that this continuous cooling transformation structure (Zw) is not easily discerned by optical microscope observation in the engraving using a nital reagent. In such a case, it is discerned by the use of EBSP-OIM ™. The Electron Back Scatter Diffraction Pattern-Orientation Image Microscopy (EBSP-OIM) method is a device and software in which a highly slanted sample on a Scanning Electron Microscope (SEM) scanning electron microscope is beamed. of electrons, a Kikuchi pattern formed by back diffusion is photographed by a high sensitivity camera, and the image is processed by a computer, and thus a crystal orientation at a point of irradiation is measured for a short period of time. time.
In the EBSP method, it is possible to quantitatively analyze a microstructure and crystal orientation of a global sample surface, and considering that an area to be analyzed is within an area capable of being observed by SEM, it is possible to analyze the area with a minimum resolution of 20 nm, depending on the resolution of the SEM. EBSP-IOM analysis is performed by mapping an area to be analyzed at ten to thousands of equally spaced grid points for several hours. You can see crystal orientation distributions and crystal grain sizes within the sample in a polycrystalline material. In the present invention, discernable from a mapped image with a packet misorientation set to 15 ° may also be defined as the continuous cooling (Zw) transformation structure for convenience.
Additionally, the structural fraction of the proeuctoid ferrite was obtained by a KAM (Kernel Average Misorientation) method being equipped with the EBSP-OIM. The KAM method is a calculation in which misorientations between six adjacent pixel pixels (first approximations) of a certain regular hex of measurement data, or 12 pixels (second approximations) positioned outside the six pixels, or 18 pixels (third approximations) additionally positioned outside the 12 pixels are averaged and a obtained value is adjusted to a central pixel value, is realized relative to each pixel.
This calculation is performed so as not to exceed a grain boundary, thereby making it possible to create a map representing a change of orientation within a grain. That is, this map represents a resistance distribution based on a local orientation change within a grain. It is noted that as the condition of analysis in the present invention, the condition of which in the EBSP-IOM, the misorientation between adjacent pixels is calculated is adjusted to the third approximation and one having this misorientation being 5 ° or less is revealed.
In the present invention, proeuctoid ferrite is defined as a microstructure up to a planar fraction of pixels of which third misorientation approximation is calculated to be 1 ° or less, as described above. This is because the high temperature polygonal pro-euctoid ferrite is generated in a diffusion transformation, and thus a displacement density is small, and resistance within the grain is small, and thus a difference within the grain. in crystal orientation is small, and, according to the results of various examinations that were performed by the present inventors, a volume fraction of polygonal ferrite obtained by optical microscope observation, and an area fraction of an area obtained by 1 ° of the third misorientation approximation measured by the KAM method substantially agree with each other. (Production Method) Hereinafter, a method of producing the cold rolled steel sheet of the present invention will be described herein. In order to achieve excellent local deformability, it is important to form a texture having predetermined pole densities to produce a steel plate that satisfies the conditions of producing fine crystal grains and equiaxial property and crystal grain homogenization. Details of production conditions to satisfy them at the same time will be described below.
A production method prior to hot rolling is not limited in particular. That is, subsequent to melting by a shaft furnace, an electric furnace, or the like, secondary refining may be variably performed, and then casting may be performed by normal continuous casting, or casting by an ingot method, or additionally. by a method such as slab casting. In the case of continuous casting, it is possible for a casting board to be immediately cooled to a low temperature and then reheated, to undergo hot rolling, or it is also possible for a casting board to be continuously hot rolled. A scrap can also be used as a raw material.
In addition, in hot rolling, it is also possible for the sheet bars to be bonded after rough rolling to be subjected to continuously finished rolling. On this occasion, it is also possible for unrolled bars to be coiled into a coil form immediately, stored in a cover having a heat insulation function as required, and rewound to be joined together. (First Hot Rolling) A board extracted from a heating furnace is subjected to a rough rolling process being the first hot rolling to be rough rolled, and thereby a raw bar is obtained. A high strength steel plate having excellent local deformability of the present invention is obtained when the following requirements are met. First, a diameter of austenite grain in the bar after gross lamination, namely before finished lamination is important, and the austenite grain diameter before finished lamination is desirably small, and it becomes clear that the diameter of Austenite grain of 200 μηη or less contributes greatly to fine grain production in the grain unit and homogenization of a main phase.
In order to obtain this austenite grain diameter of 200 μηη or less prior to finished lamination as shown in FIG. 3, In the raw rolling in a temperature region of not lower than 1000 ° C, nor higher than 1200 ° C, the rolling is performed once or more in a reduction ratio of at least 40% or more. As the reduction ratio and the number of reduction times are larger, fine grains can be obtained, and in order to efficiently achieve this effect, the austenite grain diameter is desirably adjusted to 100 μηη or less, and in order to achieve this, lamination at 40% or more is desirably performed twice or more. However, when in raw lamination, the reduction ratio is greater than 70%, and lamination is performed more than 10 times, with an interest that the temperature decreases or a scale is generated excessively.
Thus, the decrease in austenite grain diameter prior to finished lamination is effective for improving local deformability by controlling austenite recrystallization promotion in the finished lamination, producing fine grains, and producing equiaxial grains of the grain unit. in a final structure. This is supposed to be because an austenite grain boundary after gross lamination (namely, before finished lamination) functions as one of the recrystallization cores during finished lamination.
In order to confirm the austenite grain diameter after rough rolling, a sheet part previously subjected to finished lamination is desirably tempered as much as possible, and the sheet part is cooled to a cooling rate of 10 ° C / s or further, and the structure of a cross section of the sheet metal that is engraved to produce austenite grain boundaries appears, and the austenite grain boundaries are measured by an optical microscope. At this time, at 50 or more magnifications, 20 or more visual fields are measured by image analysis or a point counting method. (Second Hot Rolling) After the rough rolling process (first hot rolling) is completed, a finished rolling process being second hot rolling is started. The time between completion of the raw rolling process and the start of the finished rolling process is desirably set to 150 seconds or shorter.
In the finished lamination (second hot rolling) process, a starting temperature of the finished lamination is desirably set at 1000 ° C or higher. When the starting temperature of the finished lamination is lower than 1000 ° C, at each fitted lamination pass, the temperature of the lamination to be applied to the blank to be rolled is reduced, the reduction is performed in a region of At non-recrystallization temperature, texture develops, and thus isotropy deteriorates.
Incidentally, the upper limit of the starting temperature of the finished lamination is not limited in particular. However, when it is 1150 ° C or higher, a bubble to be a starting point for a scaly shaft-shaped scale defect is likely to occur between an iron-based steel plate and a surface scale prior to rolling. between the passages, and thus the starting temperature of the finished lamination is desirably lower than 1150 ° C.
In finished rolling, a temperature determined by the chemical composition of the steel sheet is set to T1, and in a temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C, the rolling 30% or more is performed in one pass at least once. Additionally, in finished lamination, the total reduction ratio is adjusted to 50% or more. Satisfying this condition, in the central plate thickness portion being a range of 5/8 to 3/8 plate thickness from the steel plate surface, the average value of the pole group densities {100 } <011> at {223} <110> becomes less than 4.0, and the pole density of the {332} <113> orientation group becomes 5.0 or less. This makes it possible to obtain local deformability of an end product.
Here, T1 is the temperature calculated by Expression (1) below. T1 (° C) = 850 + 10 x (C + N) x Mn + 350 x Nb + 250 x Ti + 40 x B + 10xCr + 100xMo + 100 xV ··· (1) C, N, Mn, Nb, Ti , B, Cr, Mo, and V each represent the element content (mass%).
FIGs. 4 and FIG. 5 each shows the relationship between a reduction ratio in each temperature region and a pole density of each orientation. As shown in FIG. 4 and FIG. 5, Heavy reduction in temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C, and slight reduction to T1 or higher, and lower than T1 + 30 ° C then control the mean pole density of the {100} to {223} <110> orientation group pole density, and the {332} <113> orientation group pole density at the center thickness portion of sheet being a range of 5/8 to 3/8 in sheet thickness from the steel plate surface, and thus the local deformability of the final product is dramatically improved, as shown in Tables 2 and 3 of Examples a be described later.
This temperature T1 is obtained empirically. The present inventors have learned empirically from experiments that recrystallization in an austenite region of each steel is promoted on the basis of temperature T1. In order to obtain better local deformability, it is important to build up resistance by heavy reduction, and the total reduction ratio of 50% or more is essential. Additionally, it is desired to take reduction at 70% or more, and on the other hand, if the reduction ratio greater than 90% is taken, safety temperature and excessive rolling load are, as a result, added.
When the total reduction ratio in the temperature region of not lower than T1 + 30 ° C nor higher than T1 + 200 ° C is less than 50%, the rolling resistance to be accumulated during rolling to hot is not enough, and austenite recrystallization does not advance sufficiently. Therefore, texture develops, and isotropy deteriorates. When the total reduction ratio is 70% or more, sufficient isotropy can be obtained even though variations attributable to temperature fluctuation or the like are considered. On the other hand, when the total reduction ratio exceeds 90%, it is difficult to obtain the temperature region of T1 + 200 ° C or lower due to operation heat generation, and additionally a rolling load increases to cause a risk that lamination becomes difficult to perform.
In finished lamination, in order to promote uniform recrystallization caused by release of the accumulated resistance, lamination at 30% or more is performed in one pass at least once at not less than T1 + 30 ° C, nor higher than T1 + 200 ° C.
Incidentally, in order to promote uniform recrystallization, it is necessary to suppress an amount of operation in a temperature region of as low as T1 + 30 ° C as possible. In order to achieve it, the reduction ratio lower than T1 + 30 ° C is desirably 30% or less. In terms of sheet thickness and sheet shape accuracy, the reduction ratio of 10% or less is desirable. When isotropy is additionally obtained, the reduction ratio in the temperature region of lower than T1 + 30 ° C is desirably 0%. The finished lamination is desirably finished at T1 + 30 ° C or higher. At hot rolling lower than T1 + 30 ° C, the granulated austenite grains that are recrystallized once are thereby elongated, causing a risk that isotropy will deteriorate.
That is, in the production method of the present invention, in finished lamination, by uniformly and finely recrystallizing austenite, the texture of the product is controlled and local deformability such as hole expansion or curvature is improved.
A rolling ratio can be obtained by actual performance or calculation from rolling load, sheet thickness measurement, or / and the like. Temperature can actually be measured by a thermometer between pauses, or can be obtained by calculating simulation considering heat generation by operating a line speed, reduction ratio, or / and the like. In this way, it is possible to easily confirm whether or not the lamination prescribed in the present invention is performed.
When hot rolling is finished at Ar3 or lower, the hot rolling becomes two-phase region rolling of austenite and ferrite, and buildup to the {100} To {223} <110> orientation group It becomes strong. As a result, local deformability deteriorates significantly.
In order to produce fine crystal grains and suppress elongated grains, a maximum operating heat generation amount at reduction time to not lower than T1 + 30 ° C, nor higher than Τ1 + 200 ° C, namely, an increased temperature range by reduction is desirably suppressed at 18 ° C or below. To achieve this, inter-support cooling or the like is desirably applied. (Primary Cooling) After the final reduction by a reduction ratio of 30% or more is performed on the finished lamination, primary cooling is initiated such that a wait time t second satisfies Expression (2) below. t ^ 2,5 xt1 ··· (2) Here, t1 is obtained by Expression (3) below. t1 = 0.001 x ((Tf - T1) x P1 / 100) 2 - 0.109 x ((Tf - T1) x P1 / 100) + 3.1 ··· (3) Here, in Expression (3) above, Tf represents the temperature of a steel billet obtained after the final reduction by a reduction ratio of 30% or more, and P1 represents the reduction ratio of the final reduction to 30% or more.
Incidentally, the "final reduction at a reduction ratio of 30% or more" indicates the lamination finally made between the laminations whose reduction ratio becomes 30% or more outside the laminations in a plurality of passages made in the finished lamination. For example, when between laminations in a plurality of passes made in the finished lamination, the reduction ratio of the final stage lamination is 30% or more, the final stage lamination is the "final reduction in a reduction ratio of 30% or more. " Additionally, when between laminations in a plurality of passes made in the finished lamination, the reduction ratio of the lamination performed before the final stage is 30% or more, and after the lamination performed before the final stage (lamination at a reduction ratio of 30% or more) is performed, lamination whose reduction ratio becomes 30% or more is not performed, lamination performed before the final stage (lamination at a reduction ratio of 30% or more) is the "final reduction" in a reduction ratio of 30% or more. "
In finished lamination, the wait time t second until primary cooling is initiated after the final reduction in a reduction ratio of 30% or more is performed greatly affects the austenite grain diameter. That is, it greatly affects an equiaxial grain fraction and a coarse grain area ratio of the steel plate.
When the wait time t exceeds t1 x 2.5, the recrystallization is almost complete, but the crystal grains grow significantly and the grain thickens, and thus an r value and elongation are decreased. The wait time t second additionally satisfies Expression (2a) below, thereby making it possible to preferentially suppress crystal grain growth. Accordingly, even though recrystallization does not advance sufficiently, it is possible to sufficiently improve the elongation of the steel sheet, and to improve the fatigue property simultaneously. t <t1 ··· (2a) At the same time, the wait time t second additionally satisfies Expression (2b) below, and thus recrystallization proceeds sufficiently, and the crystal orientations are randomized. Therefore, it is possible to sufficiently improve the elongation of the steel sheet, and greatly improve the isotropy simultaneously. t1 ^ t ^ t1x2,5- (2b) Here, as shown in FIG. 6, in a continuous hot rolling line 1, the steel billet (board) heated to a predetermined temperature in the heating furnace is rolled into a roughing mill 2 and a finishing mill 3 sequentially to be a sheet metal. hot-rolled steel plate 4 having a predetermined thickness, and hot-rolled steel plate 4 is made on a worktable 5. In the production method of the present invention, in the rough rolling process (first hot rolling) performed on roughing mill 2, rolling at a reduction ratio of 40% or more is performed on the steel billet (plank) once or more in the temperature range of not lower than 1000 ° C nor higher than that 1200 ° C. The blank bar is rolled to a predetermined thickness in the roughing mill 2, so it is then finely rolled (subjected to the second hot rolling) through a plurality of lamination supports 6 of the finishing mill 3 to are hot-rolled steel sheet 4. Then, in the finishing mill 3, the 30% or more rolling is performed in one pass at least once in the temperature region of not lower than the temperature T1 + 30. ° C, not higher than T1 + 200 ° C. Additionally, in finishing mill 3, the total reduction ratio becomes 50% or more.
Additionally, in the finished lamination process, after the final reduction by a reduction ratio of 30% or more is performed, primary cooling is initiated such that the wait time t second satisfies Expression (2) above or any Expression. (2a) or (2b) above. The initiation of this primary cooling is accomplished by inter-support cooling nozzles 10 disposed between the respective two of the finishing mill lamination supports 6, or cooling nozzles 11 arranged on the execution table 5.
For example, when the final reduction at a reduction ratio of 30% or more is performed only on the lamination support 6 disposed at the front stage of the finisher 3 (on the left side in FIG. 6, the upstream side of the lamination) , and the lamination whose reduction ratio becomes 30% or more is not performed on the lamination support 6 disposed at the rear stage of the finishing mill 3 (on the right side in FIG. 6, downstream side of the lamination), if the primary cooling start is accomplished by the cooling nozzles 11 arranged on the execution table 5, a case where the wait time t second does not satisfy Expression (2) above or Expressions (2a) and (2b) above is sometimes , caused. In such a case, the primary cooling is initiated by the inter-support cooling nozzles 10 disposed between the respective two of the finishing mills 6 of the finishing mill 3.
In addition, for example, when a final reduction in a reduction ratio of 30% or more is performed on the lamination support 6 disposed at the rear stage of the finishing mill 3 (on the right side in FIG. 6, downstream side of the lamination). ), even though the primary cooling start is performed by the cooling nozzles 11 arranged on the execution table 5, there is sometimes a case that the wait time t second may satisfy Expression (2) above, or Expressions (2a ) and (2b) above. In such a case, the primary cooling may also be initiated by the cooling nozzles 11 arranged on the table 5. However, assuming that the final reduction performance at a reduction ratio of 30% or more is completed, the Primary cooling may also be initiated by the inter-support cooling nozzles 10 disposed between the respective two of the lamination supports 6 of the finishing mill 3.
Then, in this primary cooling, cooling that at an average cooling rate of 50 ° C / second or more, a temperature change (temperature drop) becomes no less than 40 ° C, no more than 140 ° C. ° C, is performed.
When the temperature change is less than 40 ° C, the recrystallized austenite grains grow, and the low temperature hardness deteriorates. The temperature change is adjusted to 40 ° C or higher thereby making it possible to suppress the thickening of the austenite grains. When the temperature change is less than 40 ° C, the effect cannot be obtained. On the other hand, when the temperature change exceeds 140 ° C, recrystallization becomes insufficient to make it difficult to obtain an objective random texture. Additionally, an effective ferrite phase for elongation is also not readily obtained, and the hardness of a ferrite phase becomes high, and thus elongation and local deformability also deteriorate. Additionally, when the temperature change is greater than 140 ° C, an overrun to / beyond a transformation point temperature of Ar3 is likely to be caused. In this case, even by the transformation of recrystallized austenite as a result of variant selection alignment, the texture is formed, and the isotropy decreases accordingly.
When the average cooling rate in primary cooling is less than 50 ° C / second as expected, the recrystallized austenite grains grow and the low temperature hardness deteriorates. The upper limit of the average cooling rate is not determined in particular, but in terms of the steel sheet shape, 200 ° C / second or less is considered to be correct.
Additionally, in order to suppress grain growth, and obtain more excellent low temperature hardness, a cross-pass cooling device or the like is desirably used to bring operation heat generation between the respective finished lamination supports at 18 ° C. or lower. The rolling ratio (reduction ratio) may be obtained by actual performance or calculation from the rolling load, sheet thickness measurement, or / and the like. The temperature of the steel billet during rolling can actually be measured by a thermometer being arranged between the supports, or can be obtained by simulation by considering heat generation by operating a line speed, the reduction ratio, or / and the like, or may be obtained by both methods.
Additionally, as explained above, in order to promote uniform recrystallization, the amount of operation in the temperature region of lower than T1 + 30 ° C is desirably the lowest possible, and the reduction ratio in the temperature region of more lower than T1 + 30 ° C is desirably 30% or less. For example, in the case that in the finishing mill 3 in the continuous hot rolling line 1 shown in FIG. 6, when passing through one or two or more of the lamination brackets 6 disposed on the front stage side (on the left side in FIG. 6, on the upstream side of the lamination), the steel plate is in the temperature region of no. lower than T1 + 30 ° C, not higher than T1 + 200 ° C, and passing through one or two or more of the lamination supports 6 disposed on the subsequent rear stage side (on the right side in FIG. 6, downstream side of the rolling), the steel sheet is in the temperature region of lower than T1 + 30 ° C, when the steel sheet passes through one or two or more of the rolling supports 6 on the subsequent posterior stage side (on the right side in FIG. 6 on the downstream side of the lamination), even though the reduction is not performed, or the reduction ratio is lower than T1 + 30 ° C. it is desirably 30% or less in total. In terms of the accuracy of sheet thickness and sheet shape, the reduction ratio at lower than T1 + 30 ° C is desirably a reduction ratio of 10% or less in total. When isotropy is additionally obtained, the reduction ratio in the temperature region of lower than T1 + 30 ° C is desirably 0%.
In the production method of the present invention, a rolling speed is not limited in particular. However, when the rolling speed on the final support side of the finished lamination is less than 400 mpm, γ grains grow and thicken, the regions where ferrite can precipitate for ductility are reduced, and thus ductility. is likely to deteriorate. Even though the upper limit of the rolling speed is not limited in particular, the effect of the present invention can be obtained, but it is, in fact, that the rolling speed is 1800 mpm or less due to the restriction of ease. Therefore, in the finished lamination process, the lamination speed is desirably not less than 400 mpm, nor more than 1800 mpm.
Incidentally, after this primary cooling, coiling is performed at an appropriate temperature, and an original hot-rolled plate can be obtained. In the present invention, the cold-rolled steel sheet microstructure is mainly formed by cold rolling later, or a heat treatment after cold rolling. Thus, a cooling pattern for winding does not have to be strictly controlled much anymore. (Cold Rolling) The original hot-rolled plate produced as described above is pickled according to a need to be cold rolled at a reduction ratio of not less than 30% or more than 70%. When the reduction ratio is 30% or less, it is difficult to cause recrystallization on heating and subsequent retention resulting in the equiaxial grain fraction decreasing and additionally the crystal grains upon heating become coarse. When lamination above 70% is performed, a texture in the heating time is developed, and thus anisotropy becomes strong. Therefore, the reduction ratio is set to 70% or less. (Heating and retention) The cold rolled steel plate is then heated to a temperature region of Ae3 at 950 ° C, and is held for 1 to 300 seconds / seconds in the temperature region of Ae3 to 950 ° C. ° C to produce an austenite single phase steel or substantially an austenite single phase steel. By this heating and retention, the operation hardening is removed. In order to heat the steel plate after cold rolling to a temperature region of Ae3 at 950 ° C in this way, an average heating rate of not lower than room temperature nor higher than 650 ° C, HR1 (° C / second) expressed by Expression (5) below, and an average heating rate of higher than 650 ° C to Ae3 at 950 ° C is adjusted to HR2 (° C / second) expressed by Expression (6) below. HR1 ^ 0.3 ... (5) HR2 ^ 0.5x HR1 ... (6) Hot rolling is performed under the condition described above, and in addition primary cooling is performed, and thus producing the fine crystal grains, and the randomization of crystal orientations is achieved. However, by the cold rolling performed thereafter, the strong texture develops, and the texture becomes likely to remain in the steel plate. As a result, the value r and the elongation of the steel sheet decrease, and isotropy decreases. Thus, it is desired to make the texture developed by cold rolling disappear as much as possible by properly performing the heating to be performed after cold rolling. In order to achieve this, it is necessary to divide the average heating rate of the heating into two stages expressed by Expressions (5) and (6) above. The detailed reason why the texture and properties of the steel sheet is enhanced by this two stage heating is unclear, but this effect is thought to be related to the time-shifting recovery of cold rolling, and recrystallization. That is, the driving force of recrystallization that occurs on the steel plate upon heating is the accumulated resistance on the steel plate by cold rolling. When the average heating rate HR1 in the temperature range of not lower than room temperature nor higher than 650 ° C is small, the displacement introduced by cold rolling recovers, and recrystallization does not occur. As a result, the texture that developed at the time of cold rolling remains as it is, and properties, such as isotropy, deteriorate. When the average heating rate HR1 in the temperature range of not lower than room temperature nor higher than 650 ° C is less than 0.3 ° C / second, the displacement introduced by cold rolling recovers. resulting in the strong texture formed over time of cold rolling remains. Therefore, it is necessary to adjust the average heat rate HR1 in the temperature range of not lower than room temperature, nor higher than 650 ° C, to 0.3 (° C / second), or higher.
On the other hand, when the average heating rate HR2 higher than 650 ° C to Ae3 at 950 ° C is large, the ferrite on the steel plate after cold rolling does not recrystallize, and the non-recrystallized ferrite. in a state of being operated, remains. When the particular C-containing steel of 0.01% or more is heated to a two-phase region of ferrite and austenite, the growth of austenite blocks formed of recrystallized ferrite, and thus unrecrystallized ferrite, make become more likely to remain. This unrecrystallized ferrite has strong texture, and thus properties, such as the rea isotropy value, are adversely affected, and this unrecrystallized ferrite contains a lot of displacements, and thus ductility is deteriorated. drastically. Therefore, in the temperature range from higher than 650 ° C to Ae3 to 950 ° C, the average heat rate HR2 needs to be 0.5 x HR1 (° C / second) or less.
Additionally, at the average two-stage heat rate as above, the steel plate is heated to a temperature region of Ae3 at 950 ° C, and is maintained for 1 to 300 seconds / seconds in the temperature region of Ae3 to 950 ° C. ° C. If the temperature is lower than this range, or the time is shorter than this range, the fraction of the bainite structure does not become 95% or more in a secondary cooling process then, and the increased ductility margin. local by texture control, decreases. On the other hand, if the steel plate is continuously kept higher than 950 ° C or higher than 300 seconds, the crystal grains become coarse, and thus a grain area ratio having 20 μηη or less. , increases. Incidentally, Ae3 [° C] is calculated by Expression (7) below by the contents of C, Mn, Si, Cu, Ni, Cr, and Mo [mass%]. Incidentally, when the selected element is not contained, the calculation is performed with the selected element content [mass%] set to zero.
Ae3 = 911 - 239C - 36Mn + 40Si - 28Cu - 20Ni - 12Cr + 63Mo ... (7) Incidentally, in this heating and retention, retention not only means isothermal retention, and is sufficient if the steel plate is retained in the temperature range from Ae3 to 950 ° C. Assuming the steel sheet is in the temperature range of Ae3 to 950 ° C, the temperature of the steel sheet can be changed. (Secondary cooling) Secondary cooling is then performed at a temperature of 500 ° C or lower, so that an average cooling rate in a temperature region of Ae4 to 500 ° C can become no less than 10 ° C / s, not more than 200 ° C / s. When a secondary cooling rate is less than 10 ° C / s, ferrite is excessively generated, thereby making it impossible to bring the fraction of the bainite structure to 95% or more, and resulting in the increased local ductility margin by texture control decreases. On the other hand, even when the cooling rate is set higher than 200 ° C / s, controllability at a cooling finish temperature deteriorates significantly, and thus the cooling rate is set to 200 ° C / s. s or less. Preferably, an average cooling rate at HF (a heating and holding temperature) at 0.5HF + 250 ° C is adjusted not to exceed an average cooling rate at 0.5HF + 250 ° C at 500 ° C, from safely suppress ferrite transformation and pearlite transformation. (Excessive Heat Treatment) In order to promote bainite transformation, an excess heat treatment is carried out in a temperature range of not lower than 350 ° C nor higher than 500 ° C subsequent to cooling. secondary. A retention time in this temperature range is adjusted to t2 seconds or longer, which satisfies Expression (4) below according to an excess treatment temperature T2. However, in consideration of an applicable Expression temperature range (4), the maximum value of t2 is set to 400 seconds.
Iog (t2) = 0.0002 (T2 - 425) 2 + 1.18 ... (4) Incidentally, in this excess heat treatment, retention means not only isothermal retention, and is sufficient if the steel plate It is retained in the temperature range of not lower than 350 ° C, nor higher than 500 ° C. For example, the steel sheet may once be cooled to 350 ° C to then be heated to 500 ° C, or the steel sheet may also be cooled to 500 ° C and then cooled to below 500 ° C. 350 ° C.
Incidentally, even when a surface treatment is performed on the high strength cold rolled steel sheet of the present invention, the local deformability enhancing effect does not disappear, and for example a hot dip galvanized layer, or a layer Hot-dip galvanized alloy, can be formed on the surface of steel sheet. In this case, the effect of the present invention can be obtained even when any of electrodeposition, hot dip, deposition coating, organic coating film formation, film lamination, organic salt / inorganic salt treatment, chromium free treatment, and so on. on, it is accomplished. Additionally, the steel sheet according to the present invention may be applied not only to curl forming, but also to combined forming mainly composed of bending operation such as bending, bulging, and stretching.
Example Hereinafter, examples of the present invention will be explained. Incidentally, the conditions of the examples are condition examples employed to confirm the applicability and effects of the present invention, and the present invention is not limited to these condition examples. The present invention may employ various conditions, considering that the object of the present invention is achieved without departing from the spirit of the invention. The chemical compositions of the respective steels used in the examples are shown in Table 1. The respective production conditions are shown in Table 2 and Table 3. In addition, structural constitutions and mechanical properties of the respective steel types under the production conditions in Table 2 are shown. The structural constitutions and mechanical properties of the respective types of steel under the production conditions in Table 3 are shown in Table 5. Incidentally, each underline in the Tables indicates that a numerical value is outside the range of the present invention or outside the range of a preferred range of the present invention.
As examples, here results of examinations using steels A through T which satisfy the components of the claims of the present invention and comparative steels therein having the chemical compositions shown in Table 1 will be explained. Incidentally in Table 1 each numerical value of the Chemical compositions means mass%.
These steels were melted and then, as they were, or were reheated after once being cooled to room temperature, and were heated to a temperature range of 1000 ° C to 1300 ° C, and then subjected to hot rolling under the conditions of Table 2 and Table 3, and the hot rolling was finished at a transformation temperature of Ar 3 or higher. Incidentally, in Table 2 and Table 3, the letters A through T and English letters a through i, which are added to the steel types, indicate that they are the respective components of Steels A through T and a through i in Table 1.
In hot rolling first, in raw rolling being first hot rolling, rolling was performed once or more at a reduction ratio of 40% or more in a temperature region of not lower than 1000 ° C, nor higher than 1200 ° C. However, with respect to Steel B2, H3, and J2 types in Table 2, and Steel B2 ', H3', and J2 'types in Table 3, in rolling, rolling at a reduction ratio of 40% or more in one pass was not performed. The number of times of reduction and each reduction ratio (%) in raw rolling, and austenite grain diameter (μηη) after raw rolling (before finished rolling) are shown in Table 2 and Table 3.
After the raw rolling was completed, the finished rolling being the second hot rolling was performed. In finished lamination, lamination at a reduction ratio of 30% or more was performed in one pass at least once in a temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200. ° C, and in a temperature range of lower than T1 + 30 ° C, the total reduction ratio has been set to 30% or less. Incidentally, in finished lamination, lamination at a reduction ratio of 30% or more was performed at a final pass in the temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C. .
However, for G2, H4, and M3 Steel types in Table 2 and G2 ', H4', and M3 'Steel types in Table 3, rolling at a reduction ratio of 30% or more was not performed at temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C. Additionally, with respect to Steel types F3 and H6 in Table 2 and Steel types F3 'and H6' in Table 3, the total reduction ratio in the temperature range of lower than T1 + 30 ° C was greater than 30%
Additionally, in the finished lamination, the total reduction ratio was adjusted to 50% or more. However, for Steel types G2, H4, and M3 in Table 2 and Steel types G2 ', H4', and M3 'in Table 3, the total reduction ratio was less than 50%. Table 2 and Table 3 show, in the finished lamination, the reduction ratio (%) in the final pass in the temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C, and the reduction ratio in one pass at an earlier stage than the final pass (reduction ratio in one pass before the final) (%). In addition, Table 2 and Table 3 show, in the finished lamination, the total reduction ratio (%) in the temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C, and a temperature Tf after reduction in the final pass in the temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C. Incidentally, the reduction ratio (%) in the final pass in the temperature region of not lower than T1 + 30 ° C, nor higher than T1 + 200 ° C in the finished lamination, is particularly important in order to thereby , be shown in Table 2 and Table 3 as P1.
After the final reduction by a reduction ratio of 30% or more was performed on the finished lamination, primary cooling was initiated before a t second wait time exceeding 2.5 x t1. In primary cooling, an average cooling rate was set at 50 ° C / second or higher. Additionally, a temperature change (a cooled temperature amount) in the primary cooling has been adjusted to fall within a range of not less than 40 ° C or more than 140 ° C.
Under the production conditions shown in Table 2, after the final reduction by a reduction ratio of 30% or more was performed on the finished lamination, primary cooling was initiated before the wait time t second exceeding t1 (t <t1). On the other hand, under the production conditions shown in Table 3, after the final reduction by a reduction ratio of 30% or more was performed on the finished lamination, primary cooling was initiated before the holding time t second exceeding a range of t1 or longer at 2.5 x t1 (t1 t t1 x 2.5). Incidentally, ['] (dash) was added to each reference numeral of the steel types following the production conditions shown in Table 3 in order to distinguish the wait time ranges t second.
However, with respect to the type of steel H13 'shown in Table 3, primary cooling was initiated after the wait time second t exceeded 2.5 x t1 since the final reduction by a reduction ratio of 30% or more in finished rolling. . Regarding the type of steel M2 in Table 2 and type of steel M2 'in Table 3, the temperature change (quantity of cooled temperature) in the primary cooling was less than 40 ° C, and in relation to type H12 Steel in the In Table 2 and Steel type H12 'in Table 3, the temperature change (amount of cooled temperature) in the primary cooling was greater than 140 ° C. With respect to H8 Steel type in Table 2 and H8 'Steel type in Table 3, the average cooling rate in primary cooling was less than 50 ° C / second. Table 2 and Table 3 show t1 (second) and 2.5 x t1 (second) of the respective steel types. In addition, Table 2 and Table 3 show the wait time t (sec) for final reduction completion at a reduction ratio of 30% or more to initiate primary cooling, t / t1, the mean cooling rate (° C / second) in primary cooling, and the temperature change (cooled temperature amount) (° C).
After primary cooling, coiling was performed and the original hot-rolled sheets each having a thickness of 2 to 5 mm were obtained. Table 2 and Table 3 show the coiling temperature (° C) of the respective steel types.
The original hot-rolled sheets were then stripped and then cold rolled at a reduction ratio of not less than 30% or more than 70% at a thickness of 1.2 to 2.3 mm. However, for steel types E2 and L2 in Table 2 and steel types E2 'and L2' in Table 3, the cold rolling reduction ratio was less than 30%. Additionally, for steel grade H11 in Table 2 and steel grade H1T in Table 3, the cold rolling reduction ratio was greater than 70%. Table 2 and Table 3 show the reduction ratio (%) in cold rolling of the respective steel types.
After cold rolling, heating was performed to a temperature region of Ae3 at 950 ° C and retention was performed for 1 to 300 seconds / seconds at the temperature region of Ae3 at 950 ° C. Additionally, in order to perform heating to a temperature region of Ae3 at 950 ° C, an average heating rate HR1 (° C / second) of not lower than room temperature nor higher than 650 ° C, was adjusted to 0.3 or more (HR1 0.3), and an average HR2 heating rate (° C / second) of higher than 650 ° C to Ae3 at 950 ° C was adjusted to 0.5 x HR1 or less (HR2 → 0.5x HR1).
However, for steel types C2 and G3 in Table 2 and steel types C2 'and G3' in Table 3, a heating temperature was lower than Ae3. Additionally, with respect to steel type H10 in Table 2 and steel type H10 'in Table 3, the heating temperature was higher than 950 ° C. For steel type N2 in Table 2 and Steel type N2 'in Table 3, the retention time in the temperature region of Ae3 at 950 ° C was greater than 300 seconds. Additionally, with respect to Steel Type E2 in Table 2 and Steel Type E2 'in Table 3, the average heat rate HR1 was less than 0.3 (° C / second). With respect to Steel Types C2, H6, and H8 in Table 2 and Steel Types C2 ', H6', and H8 'in Table 3, the average heat rate HR2 (° C / second) was greater than 0, 5 x HR1. Table 2 and Table 3 show As3 (° C), heating temperature (° C), retention time (second), and average heating rates HR1 and HR2 (° C / second) of the respective steel types .
After heating and retention, secondary cooling was performed at an average cooling rate of not less than 10 ° C / s, not more than 200 ° C / s, in a temperature region of Ae3 to 500 ° C. However, for Steel Type H2 in Table 2 and Steel Type H2 'in Table 3, the average cooling rate in secondary cooling was less than 10 ° C / s. Table 2 and Table 3 show the average cooling rate (° C / second) in the secondary cooling of the respective steel types.
After secondary cooling, an excess heat treatment was performed for not less than t2 seconds, not more than 400 seconds, in a temperature region of not lower than 350 ° C or higher than 500 ° C. However, for steel type H9 in Table 2 and Steel type H9 'in Table 3, a heat treatment temperature was lower than 350 ° C, and for steel type A2 and 12 in Table 2. and Steel types A2 'and 12' in Table 3, the heat treatment temperature was higher than 500 ° C. Additionally, for Steel Type D2 in Table 2 and Steel Type D2 'in Table 3, an excess treatment time was shorter than t2 seconds, and for Steel Types A2, H9, and 12 at Table 2 and A2 ', H9', and 12 'Steel Types in Table 3, the excess treatment time was greater than 400 seconds. Table 2 and Table 3 show the heat treatment of the excess temperature, t2 (sec), and the treatment time (sec) of the respective steel types.
In all cases of Table 2 and Table 3, after excess heat treatment, 0.5% hull passage lamination was performed and material evaluation was performed. Table 4 and Table 5 show an area ratio (structural fraction) (%) of bainite, pearlite, proeuctoid ferrite, martensite, and austenite retained in a metal structure of the respective steel types. Incidentally, Table 4 shows the structural constitutions and mechanical properties of the steel types following the production conditions in Table 2. In addition, Table 5 shows the structural constitutions and mechanical properties of the steel types following the production conditions in Table 3. Incidentally, with respect to the structural fraction in Table 4 and Table 5, B means bainite, P means pearlite, F means proeuctoid ferrite, M means martensite, and rA means retained austenite. Table 4 and Table 5 show, for the respective steel types, an average pole density value of the orientation group {100} To {223} <110>, a pole density of the orientation group {332} <113>, a mean crystal grain volume diameter (unit grain size) (μηη), and a crystal grain ratio having dL / dt of 3.0 or less (equiaxed grain ratio) (% ). In addition, Table 4 and Table 5 show, for the respective steel types, tensile strength TS (MPa), elongation percentage El (%), hole expansion ratio λ (%) as an index of local deformability, and a limiting bend radius by 60 ° V-shaped bend (a sheet thickness / minimum bend radius). In a bending east, direction bending C (bending C) was performed. Incidentally, a stress test and a bend test were based on JIS Z 2241 and Z 2248 (a 90 ° V-block bend test). A hole expansion test was based on Japan Iron and Steel Federation standard JFS T1001. The pole density of each of the crystal orientations was measured using the EBSP previously described at a 0.5 μηη pitch in a 3/8 to 5/8 sheet thickness region of a cross section parallel to the rolling direction. .
As indices of hole expansion and curvature satisfying TS ^ 440 MPa, El 15%, λ 90%, and plate thickness / bending radius ^ 2.5 were adjusted as conditions. It is found that only meeting the requirements of the present invention can both have excellent bore expansion and curvature as shown in FIG. 7 and FIG. 8. LEGEND OF TABLE 3 TYPE OF STEEL; Τ1 / Ό NUMBER OF 40% REDUCTION OR MORE THAN NOT LESS THAN 1000TA OR HIGHER THAN 1220B (RINSE LAMINATION) 40% OR MORE THAN NO LESS THAN 1220Ό (OR HIGHER THREADING) GRINDING LAMINATION) AUSTENITE GRAIN DIAMETER BEFORE FINISHING ENTO / pm REDUCTION RATE TO T1 + 30 TO T1 + 200Ό (FINISHING RATE) BEFORE END TO T1 + 30 TO T1 + 200Ό (FINISHING LAMINATION) /% FINAL PASS REDUCTION REASON TO T1 + 30 TO T1 + 200Ό (FINISHING LAMINATION) /% Tf: TEMPERATURE AFTER FINAL REDUCTION TO 30% OR MORE OR MORE (FINISHING LAMINATION) /% P1 : 30% OR MORE FINAL REDUCTION REDUCTION REASON (FINISHING LAMINATION) /% t1 / s-T1X2s t: 30% FINAL LAMINATION COMPLETE TIME OR MORE FOR PRIMARY COOLING / s Vt1 - PRIMARY COOLING RATE * C / s AMOUNT OF TEMPERATURE DIMENSION IN PRIMARY COOLING FC
COOLING TEMPERATURE / COLD ROLLING RATE /% HR1 -HR2 - Αθ3 / QU HEATING TEMPERATURE / Ό RETENTION TIME / s Ae3 AVERAGE RATE AOO3 / O / s TEMP. HEAT TREATMENT / t t2 / s - RETENTION TIME / s TABLE 4 LEGEND STEEL TYPE
STRUCTURAL FRACTION AVERAGE ORIENTATION GROUP POLO DENSITIES Mean Value {100} <011> A {223} <110> POLO DENSITY {223} <113>
GRAIN UNIT SIZE / μΜ PLATE THICKNESS / MINIMUM BENDING STEEL OF THIS INVENTION COMPARATIVE STEEL TABLE 5 LEGEND TYPE OF STEEL
STRUCTURAL FRACTION GUIDANCE GROUP AVERAGE VALUE OF POLO DENSITIES {100} <011> A {223} <110> POLO DENSITY {223} <113> GRAIN UNIT SIZE / μΜ EQUIAXED GRAIN REASON
PLATE THICKNESS / BENDING THICKNESS STEEL OF THE INVENTION COMPARATIVE STEEL Code Explanation 1 continuous hot rolling line 2 roughing mill 3 finishing mill 4 hot rolled steel plate 5 work table 6 rolling mill 10 intercooling nozzle -support 11 cooling nozzle 11 CLAIMS

权利要求:
Claims (10)
[1]
1. High-strength cold-rolled steel plate having excellent local deformability, characterized in that it consists of:% by mass, C: not less than 0,02% and not more than 0,20%; Si: not less than 0,001% or more than 2,5%; Mn: not less than 0,01% and not more than 4,0%; P: Not less than 0.001% and not more than 0.15%; S: not less than 0.0005% nor more than 0.03%; Al: not less than 0.001% nor more than 2.0%; N: not less than 0.0005% nor more than 0.01%; and O: not less than 0.0005% nor more than 0.01%; where Si + Al is limited to less than 1.0%, and optionally one type or two or more types of mass%, Ti: not less than 0.001% or more than 0.20%, Nb: no less than 0.001% not more than 0.20%, V: not less than 0.001% or more than 1.0%, and W: not less than 0.001% or more than 1.0%. B: not less than 0.0001% not more than 0.0050%, Mo: not less than 0.001% not more than 1.0%, Cr: not less than 0.001% not more than 2, 0%, Cu: not less than 0,001% or more than 2,0%, Ni: not less than 0,001% or more than 2,0%, Co: not less than 0,0001% or more than 1.0%, Sn: no less than 0.0001% no more than 0.2%, Zr: no less than 0.0001% no more than 0.2%, As: no less than 0.0001% not more than 0.50% Mg: not less than 0.0001% not more than 0.010%, REM: not less than 0.0001% not more than 0.1%, Ca: not less than 0.0001% or more than 0.010%, and a remainder being composed of iron and unavoidable impurities, wherein a bainite area ratio in a metal structure is 95% or more at a central portion of sheet thickness being a range of 5/8 to 3/8 sheet thickness from the surface of the sheet steel, an average pole density value of the orientation group {100} <011> to {223} <110 > repr supported by their crystal orientations of {100}, {116} <110>, {114} <110>, {113} <110>, {112} <110>, {335} <110>, {223} <110> is 4.0 or less, and a crystal orientation pole density {332} <113> is 5.0 or less, and an average crystal grain volume diameter in the metal structure is 7 μηη or less.
[2]
High strength cold-rolled steel sheet having excellent local deformability according to claim 1, characterized in that to bainite crystal grains, a ratio of crystal grains in which a ratio of a length dL in A rolling direction for a length dt in a plate thickness direction: dL / dt is 3.0 or less is 50% or more.
[3]
High-strength cold-rolled steel sheet having excellent local deformability according to claim 1, characterized in that at the surface, a hot dip galvanized layer or a hot dip galvanized alloy layer is provided. .
[4]
Method of producing a high strength cold rolled steel sheet having excellent local deformability as defined in claim 1, characterized in that it comprises: a steel billet consisting of:% by weight, C: not less than that 0,02% or more than 0,20%; Si: not less than 0,001% or more than 2,5%; Mn: not less than 0,01% and not more than 4,0%; P: Not less than 0.001% and not more than 0.15%; S: not less than 0.0005% nor more than 0.03%; Al: not less than 0.001% nor more than 2.0%; N: not less than 0.0005% nor more than 0.01%; and O: not less than 0.0005% nor more than 0.01%; where Si + Al is limited to less than 1.0%, and optionally one type or two or more mass% types, Ti: not less than 0.001% or more than 0.20%, Nb: no less than 0.001% not more than 0.20%, V: not less than 0.001% or more than 1.0%, and W: not less than 0.001% or more than 1.0%. B: not less than 0.0001% not more than 0.0050%, Mo: not less than 0.001% not more than 1.0%, Cr: not less than 0.001% not more than 2, 0%, Cu: not less than 0,001% or more than 2,0%, Ni: not less than 0,001% or more than 2,0%, Co: not less than 0,0001% or more than 1.0%, Sn: no less than 0.0001% no more than 0.2%, Zr: no less than 0.0001% no more than 0.2%, As: no less than 0.0001% not more than 0.50%. Mg: not less than 0.0001% not more than 0.010%, REM: not less than 0.0001% not more than 0.1%, Ca: not less than 0.0001% not more than 0.010%, and the remainder being composed of iron and unavoidable impurities, performing first hot rolling in which rolling at a reduction ratio of 40% or more is performed once or more in a temperature range of not lower than 1000 ° C or higher than 1200 ° C; adjusting an austenite grain diameter to 200 μηη or less by the first hot rolling; second hot rolling in which rolling at a reduction ratio of 30% or more is carried out in one pass at least once in a temperature region of not lower than T1 + 30 ° C or higher than that T1 + 200 ° C determined by Expression (1) below; adjusting the total reduction ratio in the second hot rolling mill to 50% or more; performing final reduction at a reduction ratio of 30% or more on the second hot rolling and then initiating primary cooling such that a wait time t second satisfies Expression (2) below; adjusting an average primary cooling rate at 50 ° C / second or more, and performing primary cooling in a way that a temperature change is in a range of no lower than 40 ° C or higher than 140 ° C; performing cold rolling at a reduction ratio of not less than 30% and not more than 70%; holding for 1 to 300 seconds / seconds in a temperature region of Ae3 to 950 ° C; performing secondary cooling at an average cooling rate of not less than 10 ° C / s and no more than 200 ° C / s in a temperature region of Ae3 to 500 ° C; and performing an excess heat treatment in which the retention is performed for no shorter than t2 seconds satisfying Expression (4) below no more than 400 seconds in a temperature region of no lower than 350 ° C. not higher than 500 ° C. T1 (° C) = 850 + 10 x (C + N) x Mn + 350 x Nb + 250 x Ti + 40 x B + 10xCr + 100xMo + 100 xV ··· (1) t ^ 2.5 xt1 ··· (2) Here, t1 is obtained by Expression (3) below. t1 = 0.001 x ((Tf - T1) x P1 / 100) 2 - 0.109 x ((Tf - T1) χ P1 / 100) + 3.1 ··· (3) Here, in Expression (3) above, Tf represents the temperature of the steel billet obtained after the final reduction at a reduction ratio of 30% or more, and P1 represents the reduction ratio of the final reduction to 30% or more. Iog (t2) = 0.0002 (T2 - 425) 2 + 1.18 ... (4) Here, T2 represents an excess treatment temperature, and the maximum value of t2 is set to 400.
[5]
Method of production of high-strength cold-rolled steel sheet having excellent local deformability according to claim 4, characterized in that the total reduction ratio in a temperature range lower than T1 + 30 ° C is 30% or less.
[6]
Method of producing the high strength cold rolled steel sheet having excellent local deformability according to claim 4, characterized in that the holding time t second further satisfies Expression (2a) below. t <t1 ··· (2a)
[7]
Method of producing the high strength cold rolled steel sheet having excellent local deformability according to claim 4, characterized in that the holding time t second further satisfies Expression (2b) below. t1 ^ t ^ t1 x2,5 ··· (2b)
[8]
Method of producing the high-strength cold-rolled steel sheet having excellent local deformability according to claim 4, characterized in that the primary cooling is initiated between rolling supports.
[9]
Method of producing the high-strength cold-rolled steel plate having excellent local deformability according to claim 4, characterized in that when heating is carried out to the temperature region of Ae3 at 950 ° C after cold rolling, an average heating rate not lower than room temperature or higher than 650 ° C is adjusted to HR1 (° C / second) expressed by Expression (5) below, and an average heating rate of higher than 650 ° C the Ae3 at 950 ° C is adjusted to HR2 (° C / second) expressed by Expression (6) below. HR1 ^ 0.3 ... (5) HR2 ^ 0.5xHR1 ... (6)
[10]
Method of producing the high-strength cold-rolled steel plate having excellent local deformability according to claim 4, characterized in that it further comprises: forming a hot-dip galvanized layer or a galvanized layer of hot dipping alloy on the surface.

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EP2698440A4|2015-03-04|

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法律状态:
2018-06-05| B07A| Technical examination (opinion): publication of technical examination (opinion)|

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2018-12-04| B09A| Decision: intention to grant|

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2019-01-29| B16A| Patent or certificate of addition of invention granted|Free format text: PRAZO DE VALIDADE: 20 (VINTE) ANOS CONTADOS A PARTIR DE 12/04/2012, OBSERVADAS AS CONDICOES LEGAIS. |

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2019-11-12| B25D| Requested change of name of applicant approved|Owner name: NIPPON STEEL CORPORATION (JP) |

优先权:

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申请号 | 申请日 | 专利标题

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JP2011-089250|2011-04-13|

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JP2011089250|2011-04-13|

---

PCT/JP2012/060065|WO2012141263A1|2011-04-13|2012-04-12|High-strength cold-rolled steel sheet with excellent local formability, and manufacturing method therefor|

---