JP2008514820A - High-strength cold-rolled steel sheet with excellent shape freezing property and manufacturing method thereof - Google Patents

Abstract

The present invention provides an isotropic cold-rolled steel sheet having excellent shape freezing property and a method for producing the same.
Isotropic cold-rolled steel sheet The invention relates, in the r ₉₀ value for using the low-carbon steel obtained by adding a small amount of Ti 1.3 or less, and r _m values close to the shape fixability to 1 It is excellent and has a low Δr value of 0.15 or less, and is suitable for a steel material for an automobile outer plate in which deformation in a stretching mode mainly occurs.
[Selection] Figure 2

Description

The present invention relates to a high-strength cold-rolled steel sheet suitable for use as an outer sheet material for automobiles. More specifically, the r ₉₀ value is 1.3 or less, and the average plastic deformation ratio (r _m ) is 1. The present invention relates to a high-strength cold-rolled steel sheet excellent in shape freezing property and a method for producing the same, in which the plane anisotropy coefficient (Δr) is as low as 0.15 or less and plastic deformation isotropically occurs during press working.

  In the case of a steel sheet for automobiles, a cold-rolled steel sheet having excellent workability is required in order to perform press forming without processing defects and smoothly produce a desired shape after forming. In particular, in the case of an outer panel of an automobile, dent resistance and shape freezing property are required. Regarding the above-mentioned dent resistance, BH (Bake Hardening) steel whose strength is increased after coating has been generally used. Further, in order to improve the shape freezing property, it must be uniformly deformed in the surface direction of the plate material, and a load must be applied during press molding.

  In the case of an inner panel of an automobile, deformation of a deep drawing mode mainly occurs, so that a cold-rolled steel sheet having a large stretch ratio and a large plastic deformation ratio is advantageous when processing a plate material. In the case of a panel, deformation in a stretching mode mainly occurs, so that a cold-rolled steel sheet that is uniform in the surface direction of the plate material and has low off-axis yield strength is advantageous. Thus, using a cold-rolled steel sheet having a uniform plastic deformation in the plane direction and having a low anisotropy yield strength and an excellent shape freezing property is advantageous for manufacturing an automobile outer plate having a complicated shape.

The draw ratio is a value that indicates the properties of the steel sheet that is stretched without cracking during tension. The greater the stretch ratio, the greater the allowable deformation of the steel sheet, and the stretch ratio is the steel that does not change significantly once the steel type is determined. It is a mechanical property. The plastic deformation ratio r value is a value defined by the ratio of the deformation rate in the width direction to the deformation rate in the thickness direction. A steel sheet with a large plastic deformation ratio assumes that the amount of deformation in the width direction is constant, and when the plate material is pulled in an arbitrary direction by a certain amount of deformation, the deformation rate in the thickness direction is small, and even when the amount of deformation is large, material necking (Necking) does not occur and it means that processing is possible. The plastic deformation ratio has different values along the tensile direction due to the anisotropy of the plate material. As indicating the variation degree of the plastic deformation ratio by pulling direction, there is an average plastic deformation ratio r _m and the planar anisotropy coefficient Δr value, the value can be calculated from the following equation.

r _m = (r ₀ + 2r ₄₅ + r ₉₀ ) / 4
Δr = (r ₀ −2r ₄₅ + r ₉₀ ) / 2
Here, r ₀ , r ₄₅ , and r ₉₀ mean values of plastic deformation ratios when the tensile directions are 0 °, 45 °, and 90 ° directions with respect to the rolling direction of the plate material, respectively.

When the value of the plastic deformation ratio is large, the off-axis yield strength increases and it becomes difficult to process an external panel having a complicated shape. FIG. 1 shows a theoretical analysis of the effect of the plastic deformation ratio on the yield locus of steel with two different textures, using a Taylor polycrystal model. This is the predicted result. In the case of IF steel (Interstitial Free steel) with a large plastic deformation ratio, even if the yield strength in the rolling direction (RD, rolling direction) is the same as that of an isotropic steel (isotropy steel) having a value of r _m = 1, It can be seen that the yield strength value has a large value. Thus, attempts to reduce the different shafts yield strength, it is preferable that r _m value is smaller to be close to 1. In order to reduce the Δr value, the difference in plastic deformation ratio between the directions must be small when pulled in each direction. That is, a small Δr value means that the distribution of the deformation rate in the surface direction of the plate material is uniform during press forming, which is advantageous for forming while inducing uniform deformation in the deformation in the stretching mode. It is. Thus steel having a r _m value and a low Δr value close to 1, so improving the shape fixability of the steel when the automobile outer panel machining stretching mode deformation occurs primarily.

  It is as follows if the well-known technique regarding improving the formability of the steel plate for motor vehicles is seen.

In Patent Document 1, in order to improve the formability of an automotive steel sheet, Ti or Nb is added singly or in combination to an ultra-low carbon cold-rolled steel sheet, and solid solution C and N are precipitated in the form of carbides and nitrides, and the draw ratio is increased. And the technique which improves a moldability by raising a plastic deformation ratio is proposed. Patent Documents 2 to 6 propose techniques for reducing in-plane anisotropy and reducing defects such as surface defects during press working to improve formability. The above conventional technique is a technique for reducing the in-plane anisotropy of an ultra-low carbon cold-rolled steel sheet by reducing the size of crystal grains of the hot-rolled structure immediately after finish rolling through rapid cooling equipment. Ti, ultra low carbon steel with added Nb exhibits excellent workability in deformation of the relatively high deep drawing mode r _m value and Δr value, but severe anisotropy is a variant of the stretching mode, different shafts The yield strength is high, which is disadvantageous in terms of shape freezing. Further, the above conventional technique generally cannot add high strength because only 0.005% or less of carbon is added due to deep workability.

On the other hand, Patent Documents 7 to 9 have isotropic plasticity by adding carbide forming elements such as Ti, Nb, and V in ultra-low carbon steel to control carbides and fine texture during hot rolling and annealing. A method for producing a high-strength cold-rolled steel sheet is presented. However, the conventional technique has a problem that productivity per unit time is low because a batch annealing (BAF) equipment that requires a long time is used. Further, Patent Document 10 presents a technique for producing a high-strength isotropic steel using low-carbon steel added with Ti or Nb by using a continuous annealing facility. This technique is applied to a steel sheet added with Ti or Nb. The purpose was to produce a steel sheet having a low Δr value by utilizing the strong correlation between the recrystallization extension and the Δr value. That is, the above conventional art in the production of automobile parts to be molded in a circular or rectangular tube, although manufactured Δr value to reduce the ear generation amount less than 0.1, r _m value is 1 They did not recognize that they had excellent shape freezing properties when approaching. Further, in the case of steel for automobile outer plates, the aging index (AI) is generally required to be 30 MPa or less. However, the conventional technique is applied to a steel plate formed of a cylinder or a square tube. For this reason, I was not interested in the aging index.

  On the other hand, Patent Document 11 proposes a technique for producing a high-strength isotropic steel sheet using low-carbon steel added with Ti and / or Nb using a continuous annealing facility. Since it is for producing non-aging low carbon which does not require treatment, at least one component of Cu, V and Ni other than Ti and Nb must be added at a maximum of 0.15%. Further, the steel plate of the above prior art has an isotropic value because the Δr value has a value in the range of 0.15 to 0.28.

Japanese Published Patent Publication No. 9-296226 Japanese Patent Publication No. Hei 6-158176 Japanese Published Patent Publication No. 8-109416 Japanese Patent Publication No. 11-40531 Japanese Patent Publication No. Hei 4-95392 Japanese Published Patent Publication 2002-3951 German published patent publication DE3843732 German Published Patent Publication DE3803064 US Published Patent Publication US 5,139,580 Japanese Published Patent Publication No. Hei 10-130780 US Published Patent Publication US 6,162,308

The present invention is intended to solve the problems of the conventional technology described above, with the r ₉₀ value for using the low-carbon steel obtained by adding a small amount of Ti 1.3 or less, and r _m value is 1 A high-strength isotropic cold-rolled steel sheet excellent in shape freezing property and excellent in shape freezing property, and excellent in shape freezing property, and having a Δr value as low as 0.15 or less and mainly causing deformation in a stretching mode, and its Its purpose is to provide a manufacturing method.

In the present invention, C is 0.01 to 0.05%, Ti is 0.005 to 0.06%, Mn is 0.1 to 1%, Si is 0.1% or less, and P is 0% by weight. 0.03% or less, S is 0.08% or less, N is 0.01% or less, N is 0.01% or less, the balance is Fe and inevitable impurities. Of the above components, Ti and N are Ti / N> 5, and among the above components, Ti and C are (48/12) C-Ti ^* > 0.03% [where Ti ^* = Ti- (48/14) N] And a high-strength cold-rolled steel sheet excellent in shape freezing property having an aging index (AI) of 30 MPa or less.

  The C content of the cold-rolled steel sheet is more preferably 0.015 to 0.035%.

  The Ti content of the cold-rolled steel sheet is more preferably 0.01 to 0.04%.

The cold-rolled steel sheet more preferably limits (48/12) C—Ti ^* to 0.06 to 0.11%.

Moreover, this invention is weight%, C is 0.01-0.05%, Ti is 0.005-0.06%, Mn is 0.1-1%, Si is 0.1% or less, P is 0.03% or less, S is 0.03% or less, soluble Al is 0.08% or less, N is 0.01% or less, the balance is Fe and inevitable impurities, and Ti and N in the above components are Ti / N> 5 is satisfied, and Ti and C in the above components satisfy the relationship of (48/12) C—Ti ^* > 0.03% [where Ti ^* = Ti− (48/14) N], Finishing and rolling a steel having an aging index (AI) of 30 MPa or less to a temperature of Ar ₃ or higher, and after rapidly quenching the hot-rolled steel sheet that has been finish-rolled at a cooling rate of 50 ° C./sec or higher, a temperature of 650 ° C. or lower. A step of winding at 50%, and a step of cold rolling at a rolling reduction of 50 to 80% after pickling the wound hot rolled sheet The cold-rolled cold-rolled sheet was heated at a temperature of recrystallization temperature to Ac _3, the steps of annealing, and primary cooling thereafter to 600~700 ℃ 3 ℃ / sec or more cooling rate, again 100 Secondary cooling at a cooling rate of 30 ° C./sec or more to ˜500 ° C., over-aging treatment of the cooled steel plate at a temperature of 200 to 500 ° C. within 10 minutes, and the over-aging steel plate of 0 And a step of tempering at a rolling reduction of 5% or more, and a method for producing a high-strength cold-rolled steel sheet having excellent shape freezeability.

  In the manufacturing method of the said cold-rolled steel plate, it is more preferable that C content is 0.015-0.035%.

  In the manufacturing method of the said cold-rolled steel plate, it is more preferable that Ti content is 0.01 to 0.04%.

In the manufacturing method of the said cold-rolled steel plate, it is more preferable to limit (48/12) C-Ti ^* to 0.06-0.11%.

  In the manufacturing method of the cold-rolled steel sheet, it is more preferable that the rapid cooling is performed within 1 second after finishing the finish rolling when the rapid cooling is performed after the finish rolling.

  In the manufacturing method of the said cold-rolled steel plate, it is preferable to perform the said annealing at 760-820 degreeC for 5 minutes or less.

  In the manufacturing method of the said cold-rolled steel plate, it is more preferable to restrict | limit a temperature increase rate to 3 degrees C / sec or more in the case of the temperature increase for the said annealing.

According to the present invention, r ₉₀ value is 1.3 or less, and r _m values superior near shape fixability to 1, and deformation automotive exterior primarily occurring in the lower stretching mode Δr value is 0.15 or less It is possible to provide a high-strength isotropic cold-rolled steel sheet excellent in shape freezing property suitable for a steel material for a plate and a method for producing the same. The steel according to the present invention as described above can easily process a part having a complicated shape in a stretching deformation mode when forming an automobile part.

The present inventors have found that when r _m values as low as close to 1, that is as different shafts yield strength superior becomes shape fixability depressed theoretically apparent, low-carbon was added only a small amount of Ti Research was repeated on a technique for producing a cold-rolled steel sheet having excellent shape freezing property, isotropic property, and an aging index of 30 MPa or less so as to be suitable for use in an automobile outer plate material using steel. As a result, by adjusting the components so as to have a certain relationship between Ti, N, and C, and appropriately controlling the steel sheet production conditions, particularly hot rolling conditions and annealing conditions, the r ₉₀ value is 1.3. hereinafter, and r _m value not only close to 1, found that Δr value can be aging index as low as 0.15 or less to produce steel plate of high strength continuous annealing equipment is below 30 MPa, the present invention It came to.

  First, the chemical components of the steel sheet of the present invention and the reasons for limitation will be described.

C in steel exists as an interstitial solid solution element in the form of cementite, but has a great influence on the strength of steel sheets and the formation of texture in the cold rolling and annealing processes. In the present invention, it is preferable to limit the C content in steel to 0.01 to 0.05% for the following reason. If the C content is less than 0.01%, the strength decreases and the Δr value becomes too large. Therefore, the C content is preferably limited to 0.01% or more. Moreover, since the above C combines with Fe to form cementite, it can exist stably in the steel. According to the present invention, it was revealed that an appropriate amount of C exists and needs to be precipitated with cementite in order to suppress aging at room temperature. When the amount of C is excessively large, the strength is greatly increased, the ductility is decreased, and the cold rolling property is deteriorated. Therefore, the maximum content is preferably limited to 0.05% or less. More preferably, the C content is limited to 0.015 to 0.035%. When heating in the continuous annealing process, C contained in the steel combines with Ti to precipitate TiC, resulting in an increase in strength due to the effect of precipitation strengthening, and ND (Normal Direction) which is advantageous for reducing the Δr value. Plays a role in slowing the recovery and recrystallization speed of crystal grains having an orientation parallel to the <111> direction (<111> // ND), and as a result, the fraction of crystal grains having the <111> // ND orientation. The rate is lowered. At this time, a very part of C is precipitated by Ti ₄ C ₂ S ₂ at a high temperature, but the size of the precipitate is relatively coarse compared to TiC, and the development of the recrystallization ratio orientation is almost affected. Does not reach.

  Ti, together with C, is one of the most important elements in the present invention. The Ti has an effect of suppressing the formation of AlN by combining with N as well as C to form TiN. AlN formed during hot rolling has a problem of extending the hot rolled structure and increasing the shape anisotropy of the plate. Thus, Ti suppresses the formation of AlN and reduces the fraction of crystal grains that precipitate TiC and have a strong anisotropy, thereby reducing the Δr value and increasing the strength by precipitation hardening. is there. However, since Ti is expensive, it is advantageous from the economical aspect to use as small an amount as possible. Therefore, in the present invention, it is preferable to limit the Ti content to 0.005 to 0.06% in consideration of economics while obtaining the effect of addition of Ti. A more preferable range of Ti is 0.01 to 0.04%. At this time, in order to suppress the formation of AlN and allow TiC to precipitate during annealing, the Ti / N ratio should be added so that Ti / N> 5 as much as possible. Moreover, in order to suppress normal temperature aging, an appropriate amount of C must be present and precipitated with cementite. Therefore, in the present invention, the Ti content is a function of C and N, and is limited as shown in the following formula. Is preferred.

(48/12) C-Ti ^* > 0.03%, Ti ^* = Ti- (48/14) N
Here, Ti ^* is an effective Ti content, and refers to the amount of Ti necessary for TiC formation excluding the amount of Ti necessary to form TiN in order to suppress the formation of AlN during the hot rolling process. A more preferable range of (48/12) C-Ti ^* , which is the ratio of the effective Ti (Ti ^* ) and C, is 0.06 to 0.11%.

  Mn in steel is an element effective for the effect of solid solution strengthening. In particular, S in steel is precipitated at high temperature with MnS, and during hot rolling, the occurrence of plate breakage due to S and high temperature embrittlement are suppressed. According to the experiment relating to the present invention, when the Mn content is less than 0.1%, the effect of increasing the strength cannot be obtained, and S in the steel cannot be completely precipitated with Mn. Was shown to have a problem. On the other hand, if it exceeds 1%, the effect increase due to the addition is saturated.

  In the steel, Si acts as a solid solution strengthening element. In the present invention, Si is preferably limited to 0.1% or less in order to ensure a suitable stretch ratio.

  The higher the content of P in steel, the more advantageous it is to increase the strength. However, excessive P addition increases the possibility of brittle collapse and increases the possibility of plate breakage of the slab during hot rolling, facilitating diffusion and segregation to grain boundaries after annealing is completed. Problems associated with the occurrence of secondary processing brittleness during molding are increased. Therefore, it is necessary to use it with its content limited. In the present invention, the necessary strength can be ensured by the effect of precipitation strengthening by TiC. Therefore, the P content is preferably limited to 0.03% or less.

  Since S and N are elements inevitably contained as impurities in the steel, it is important to manage the content as low as possible. However, in order to manage the content as low as possible, the refining cost of the steel increases. There is a problem. Therefore, it is preferable to manage the content as low as possible within the operating conditions. In the present invention, the S content is preferably limited to 0.03% or less. Also, since the N content changes the effective Ti content that forms TiN at high temperature and combines with C, if the N content is increased, the effective Ti content decreases. Therefore, in the present invention, it is preferable to limit the N content to 0.01% or less.

  Soluble Al effectively acts as a deoxidizing element for molten steel, but if too much Al is added, the workability may be adversely affected, so it is preferable to limit its content to 0.08% or less. .

  Next, the production method of the present invention will be described in detail.

In the present invention, the base material of the hot-rolled steel sheet may be used as it is without being produced by continuously casting a steel having the above composition range with an ingot, or may be used after being reheated after being produced with an ingot. . However, when it is intended to be reheated after being manufactured with an ingot, it is preferable to reheat the Ti ₄ C ₂ S ₂ formed at the time of cooling at 1200 ° C. or higher by cooling with the ingot.

The hot rolling process of the present invention is preferably performed according to a normal process, and is finished in a temperature range where the final pass temperature of finishing mill is Ar ₃ or higher. When the above hot rolling end temperature is lowered, the surface layer and edge portion of the hot rolled sheet are rolled in an abnormal region, the size of the crystal grains becomes coarse and becomes non-uniform, and surface defects of the material occur during press forming. It will cause. Cooling by ROT (Run Out Table) after finish rolling must be rapidly cooled to the coiling temperature at a rate of 50 ° C./sec or more to reduce the size of crystal grains of the hot rolled sheet. If the cooling rate is less than 50 ° C./sec, the crystal grains become coarse. In the above cooling, it is more preferable that the crystal grains are made finer if they are rapidly cooled within 1 second after finishing rolling. The rapid cooling can be performed using a high density cooler provided at the front end of the ROT. The winding temperature after the finish rolling is preferably limited to 650 ° C. or less. The reason for this is that when the coiling temperature exceeds 650 ° C., the TiC precipitates become coarse, and the role of slowing the recovery and recrystallization speed of subcrystal grains having a strong anisotropy during annealing becomes anisotropic. This is because the fraction of crystal grains having a strong orientation is high.

  The hot-rolled sheet wound as described above is preferably subjected to cold rolling at a rolling reduction of 50 to 80% after pickling using a normal process. When the cold rolling reduction is less than 50%, the annealing does not sufficiently recrystallize and the ductility decreases, and when it exceeds 80%, the in-plane anisotropy increases.

The annealing in the present invention is based on continuous annealing as shown in FIG. 2, and is performed in a temperature range from the recrystallization temperature to Ac ₃ points. When the annealing temperature exceeds the Ac ₃ point, annealing is performed in the abnormal region of α and γ, so that recrystallization of crystal grains having a strong anisotropy and growth of crystal grains are accelerated, resulting in coarsening of the crystal grains. become. In this way, when the crystal grains become coarse, the strength and ductility are deteriorated together. Therefore, the annealing temperature is preferably limited to Ac ₃ points or less. Moreover, when the said annealing temperature is too low with less than recrystallization temperature, there exists a problem which ductility falls. More preferably, the preferable annealing temperature is set to 760 to 820 ° C. Moreover, it is preferable that the maintenance time at the annealing temperature is within 5 minutes. The reason for this is that when annealing is performed, the longer the maintenance time, the larger the r ₉₀ value, leading to the growth of crystal orientation with large anisotropy.

  Moreover, it is preferable to perform the temperature increase to the annealing temperature after cold rolling at a temperature increase rate of 3 ° C./sec or more. This is because if the temperature rising rate is less than 3 ° C./sec, the annealing time becomes long and the recrystallized grains may become coarse.

  Thereafter, C is primary cooled to 600 to 700 ° C. where the solid solubility in the Fe matrix is high, and immediately cooled to 100 to 500 ° C. where C is low in solid solubility, so that cementite is separated from the grain boundaries and It will lead to precipitate at the interface. The primary cooling is preferably performed at a cooling rate of 3 ° C./sec or more. However, when cooling at less than 3 ° C./sec, the supersaturated C cannot be sufficiently precipitated as cementite and is ductile. In addition, the normal temperature aging deteriorates. When the secondary cooling is completed as described above, an overaging treatment is performed for 10 minutes or less so that the cementite grown by reheating at a temperature in the range of 200 to 500 ° C. can be grown. If the overaging temperature is less than 200 ° C., cementite cannot be grown sufficiently. Therefore, part of C is dissolved in solid and ductility and normal temperature aging are deteriorated. There is a problem that ductility and normal temperature aging deteriorate.

  Thereafter, the steel sheet that has been overaged as described above is subjected to temper rolling, and at this time, the rolling reduction is preferably limited to 0.5% or more.

  The following examples are provided for a more complete understanding of the invention. The specific techniques, conditions, materials, fractions and reported data presented to illustrate the principles and practices of the present invention are illustrative and do not limit the scope of the invention.

  The Ti-added carbon steel having the components shown in Table 1 below was melted and continuously cast, then reheated at 1200 ° C. and finish-rolled to 2.5 mm at a finish rolling temperature of 870 to 890 ° C. Then, after the cooling start time of the following Table 2 passed using the high-density cooling device, after rapid cooling at a rate of 60 ° C./sec, the film was wound at the winding temperature of the following Table 2. Then, after removing the surface oxide layer of the hot-rolled steel sheet by pickling, cold rolling was performed at a reduction rate of 70% up to 0.75 mm. Thereafter, the cold-rolled steel sheet was heat-treated in a continuous annealing line. The maximum heating temperature during the heat treatment was 780 to 800 ° C. After heating at the above temperature for 1 minute, primary cooling was performed to 700 ° C. at a cooling rate of 5 ° C./sec, and then secondary cooling was performed to 100 ° C. at a cooling rate of 60 ° C./sec. Then, after reheating at 300 to 350 ° C. and overaging for 3 minutes, temper rolling was performed at a rolling reduction of 1 to 1.3%. The tensile test of the annealed plate obtained as described above was performed by processing EN10002-1 test pieces.

Table 2 below shows production conditions for producing cold-rolled steel sheets having the components shown in Table 1 and the results of uniaxial experiments. In Table 2, FDT is the finish rolling finish temperature, CT is the coiling temperature, ST is the annealing temperature, YP is the yield strength, TS is the tensile strength, El is the total stretch ratio, r ₉₀ is the plasticity in the 90 ° direction from the rolling direction. Deformation ratio, Δr is a plane anisotropy coefficient, and AI is an aging index. Here, AI was calculated using the difference between the flow stress and the flow stress after heating at 100 ° C. for 1 hour after adding 7.5% pre-strain before heating.

The steels A to H in Table 2 are low steels satisfying the components of the present invention and the production conditions thereof, with an r ₉₀ value of 1.3 or less and a Δr value of 0.15 or less and low off-axis yield strength. It can also be seen that the steel has low in-plane anisotropy. On the other hand, steels I to L are steels outside the scope of the present invention, and the Ti addition amount is lower than the added N amount. That is, it was shown that the Ti / N ratio is lower than 5 which is the range of the present invention, and the Δr value is 0.15 or more. In particular, comparative steels I and J have a cooling start time longer than the limit of the present invention after finish rolling. In comparison steels M and N, although the Ti / N ratio was within the range of the present invention, r ₉₀ , Δr, and aging index could not satisfy the range of the present invention. This is because in the case of steel M, the coiling temperature is higher than the range of the present invention, and solid solution C precipitates TiC in the hot-rolled sheet and becomes coarse, thereby causing insufficient precipitation during annealing. ₉₀ , it is judged that isotropic steel cannot be obtained by increasing the development of crystal orientation ({554} <225>) having a large Δr. In the case of steels M and N, the Ti content does not satisfy (48/12) C—Ti ^* ≧ 0.6, which is a more preferable conditional expression of the present invention, and it is judged that the aging index is high. Steel O also has a cooling start time that is outside the scope of the present invention, which is coarser than when the hot-rolled sheet structure has a faster cooling start time, and after annealing, the place of nucleation of cementite during cooling is reduced. Thus, it is determined that the room temperature aging value has a high value and Δr has a value of 0.15 or more.

FIG. 3 shows a Ψ ₂ = 45 ° cross section of the orientation distribution function for the main texture component developed in steel. FIG. 4 shows the result of theoretical calculation of the influence of the texture on the anisotropy of the plastic deformation ratio of the main texture component shown in FIG. 3 using Taylor polycrystal theory. . It can be seen from FIG. 4 that α-fibre (RD // <110>) and γ-fibre (ND // <111>) textures have different effects on the plastic deformation ratio. The α-fibre texture has an overall low plastic deformation ratio and r ₄₅ has the maximum value, whereas the {554} <225> texture including γ-fibre has the lowest value of r ₄₅ . Thus, it can be seen that in order for steel to be isotropic, the above textures must be appropriately combined.

FIG. 5 shows a crystallographic orientation map (COM, Crystallographic Orientation Map) measured using EBSD (Electron Back Scattered Diffraction) equipment attached to FE (Field Emission) -SEM of Steel A of the present invention. Comparing the colors on the inverse pole figure at the top, it can be seen that α-fibre and γ-fibre textures have developed together in the steel of the present invention. FIG. 6 shows the result of analyzing not only crystal grains but also cementite together using an optical microscope. It can be seen that cementite is mainly located at the grain boundaries. Fig. 7 shows the correction of the pole figure data obtained by using X-ray diffraction for the microtexture developed in the invention steel A, and the orientation distribution function (ODF, Orientation Distribution Function) of the steel is calculated using the data. The result is shown as ψ ₂ = 45 °. FIG. 7 shows that α-fibre and γ-fibre textures are developed together. From the above, it can be seen that the steel of the present invention has excellent isotropy because the α-fiber and γ-fire textures develop together.

  Although the present invention has been described with reference to the posted embodiments, those having ordinary skill in the art will appreciate that the specific embodiments described above are merely illustrative of the invention. It can be understood that various changes can be made without departing from the spirit of the present invention.

It is a figure which shows the relationship between a plastic deformation ratio and a yield surface. It is a figure which shows the continuous annealing process by this invention, and the change of the fine structure by it. It is a figure which shows the component of the main texture which develops in steel. It is a figure which shows the influence of the texture which acts on r value. It is the crystallographic orientation figure measured using EBSD after carrying out continuous annealing of this invention steel A. It is a photograph of the optical microscope obtained after carrying out continuous annealing of this invention steel A. It is sectional drawing in (PSI) ₂ = 45 degree of the orientation distribution function measured after continuous annealing of this invention steel A. FIG.

Claims (11)

By weight, C: 0.01 to 0.05%, Ti: 0.005 to 0.06%, Mn: 0.1 to 1%, Si: 0.1% or less, P: 0.03% or less S: 0.03% or less, soluble Al: 0.08% or less, N: 0.01% or less, balance Fe and unavoidable impurities,
Among the components, Ti and N satisfy the relationship of Ti / N> 5,
Further, among the above components, Ti and C satisfy the relationship of (48/12) C-Ti ^* > 0.03% [where Ti ^* = Ti- (48/14) N],
A high-strength cold-rolled steel sheet excellent in shape freezing property having an aging index (AI) of 30 MPa or less.   The high-strength cold-rolled steel sheet having excellent shape freezing property according to claim 1, wherein the C content is 0.015 to 0.035%.   The high-strength cold-rolled steel sheet having excellent shape freezing properties according to claim 1, wherein the Ti content is 0.01 to 0.04%. The high strength cold rolling excellent in shape freezing property according to any one of claims 1 to 3, wherein the (48/12) C-Ti ^* is 0.06 to 0.11%. steel sheet. By weight, C: 0.01 to 0.05%, Ti: 0.005 to 0.06%, Mn: 0.1 to 1%, Si: 0.1% or less, P: 0.03% or less S: 0.03% or less, soluble Al: 0.08% or less, N: 0.01% or less, balance Fe and unavoidable impurities,
Among the components, Ti and N satisfy the relationship of Ti / N> 5,
Further, among the above components, Ti and C satisfy the relationship of (48/12) C-Ti ^* > 0.03% [where Ti ^* = Ti- (48/14) N],
Finishing and rolling a steel having an aging index (AI) of 30 MPa or less at a temperature of Ar ₃ or more;
The finish-rolled hot-rolled sheet is rapidly cooled at a cooling rate of 50 ° C./sec or more and then wound at a temperature of 650 ° C. or less;
After pickling the wound hot-rolled sheet, cold rolling at a rolling reduction of 50 to 80%;
Heating the cold-rolled cold-rolled sheet at a recrystallization temperature to a temperature of Ac ₃ and annealing,
Thereafter, primary cooling to 600 to 700 ° C. at a cooling rate of 3 ° C./sec or more, and secondary cooling to 100 to 500 ° C. at a cooling rate of 30 ° C./sec or more,
A step of over-aging the cooled steel sheet at a temperature of 200 to 500 ° C. for a time of 10 minutes or less;
Refining the over-aged steel sheet at a rolling reduction of 0.5% or more;
A method for producing a high-strength cold-rolled steel sheet having excellent shape freezing properties.   The method for producing a high-strength cold-rolled steel sheet having excellent shape freezing property according to claim 5, wherein the C content is 0.015 to 0.035%.   The method for producing a high-strength cold-rolled steel sheet having excellent shape freezing properties according to claim 5, wherein the Ti content is 0.01 to 0.04%. 8. The high-strength cooling with excellent shape freezing property according to claim 5, wherein the (48/12) C—Ti ^* is 0.06 to 0.11%. A method for producing rolled steel sheets.   6. The method for producing a high-strength cold-rolled steel sheet having excellent shape freezing property according to claim 5, wherein the rapid cooling is performed within 1 second after the finish rolling is finished.   The method for producing a high-strength cold-rolled steel sheet having excellent shape freezing properties according to claim 5, wherein the annealing is performed at 760 to 820 ° C for 5 minutes or less.   The method for producing a high-strength cold-rolled steel sheet having excellent shape freezing property according to claim 5 or 10, wherein a rate of temperature rise is 3 ° C / sec or more at the time of temperature rise for the annealing. .

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