JP4959418B2 - High-strength cold-rolled steel sheet and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet suitable for automobile applications and the like and a method for producing the same.

  In order to reduce the amount of carbon dioxide emissions, the weight reduction of automobile bodies for the purpose of improving the fuel efficiency of automobiles is being promoted. Therefore, the application of high-strength steel sheets capable of reducing the plate thickness is increasing for automobile members. Further, in order to ensure the safety of passengers, high-strength steel plates are increasingly used in automobile bodies.

  On the other hand, in order to apply high-strength steel sheets to automobile bodies, excellent workability is also required. In addition to improving strength and workability, specifically, yield ratio obtained by dividing yield strength by tensile strength. There is a need for reduced steel sheets. Lowering the yield ratio, that is, lowering the yield ratio has the effects of improving the shape freezing property, which is worsened by increasing the strength, reducing the press load, and suppressing the generation of wrinkles.

  Up to now, in order to achieve both low yield ratio and high strength of steel sheet, the metal structure is a composite structure of ferrite and hard second phase, specifically martensite, bainite and retained austenite. However, for example, Patent Documents 1 to 3 have proposed. However, although the cold-rolled steel sheets proposed in Patent Documents 1 to 3 have a low yield ratio, they do not consider a texture that has a very high correlation with the shape freezing property.

  On the other hand, in order to enhance the shape freezing property, a method for producing a cold-rolled steel sheet in which the development of the {554} <225>, {111} <112> and {111} <110> orientations is suppressed is disclosed in Patent Document 4, for example. Proposed. However, although the cold-rolled steel sheet proposed in Patent Document 4 is excellent in shape freezing property, the low yield ratio is insufficient.

  On the other hand, a cold-rolled steel sheet in which the metal structure is a composite structure of ferrite and a hard second phase, and further, the development of {554} <225>, {111} <112> and {111} <110> orientations is suppressed. However, it is proposed by patent document 5, for example. However, even with the cold-rolled steel sheet proposed in Patent Document 5, the yield ratio cannot be 0.55 or less.

  Patent Document 6 proposes a high-strength cold-rolled steel sheet composed of non-recrystallized ferrite and a hard second phase. However, although the strength is high, the yield ratio is high and the elongation is low. Was insufficient.

  Furthermore, in Patent Document 7, among the {001} <011> to {223} <110> orientation groups, in particular, the {112} <110> orientation and the {001} <110> orientation are developed to produce fine precipitation. Steel sheets in which objects are dispersed have been proposed. However, although the steel sheet proposed in Patent Document 7 has good shape freezing property and stretch flangeability, the yield ratio is high.

  Incidentally, it is effective to increase the yield strength in order to improve the collision resistance. However, as described above, it is effective to reduce the yield strength in order to improve the shape freezing property. For this reason, in order to achieve both the impact resistance and the shape freezing property, it is extremely important to increase the yield strength of the member by the paint baking process after forming, and the paint bake hardenability (also called Bake Hardness: BH property) of the steel sheet. It is desirable to improve.

  As such a steel sheet that achieves both shape freezing property and paint bake hardenability, the {001} <011> to {223} <110> orientation groups are developed, and the {554} <225> orientation, {111} Patent Document 8 proposes a cold-rolled steel sheet in which the development of <112> orientation and {111} <110> orientation is suppressed. However, the steel sheet proposed in Patent Document 8 is slightly inferior in forming margin during bending.

  Further, some of the present inventors include Mo or W and further B in Patent Document 9, and {001} <011> to {223} <110> orientation group, particularly {100} <011>. A steel plate that has been developed to improve shape freezeability by suppressing the development of the {554} <225>, {111} <112>, and {111} <110> orientations, and further, solid solution N. We proposed a steel plate that was used to achieve both paint bake hardenability. However, the steel sheet proposed in Patent Document 9 does not utilize unrecrystallized ferrite.

JP-A-57-2840 JP-A-62-74024 Japanese Patent Laid-Open No. 2001-220461 JP 2004-131771 A JP 2005-256020 A Japanese Patent Laid-Open No. 53-5018 JP 2006-22349 A JP 2005-12059 A Japanese Patent Application No. 2006-340069

  The present invention has a high strength cold rolling with a tensile strength of 590 MPa or more, a yield ratio of preferably 0.55 or less, excellent shape freezing properties, and more preferably also a paint bake hardenability. It is an object of the present invention to provide a steel plate and a method for manufacturing the steel plate.

The present inventors added Nb and Ti, which are elements that suppress recrystallization, optimized the cold rolling rate and annealing conditions, and had a texture that has a texture that suppresses the growth of crystal orientation that deteriorates shape freezeability. We examined the active use of crystalline ferrite and succeeded in increasing the strength and reducing the yield ratio of cold-rolled steel sheets without impairing the shape freezeability. Furthermore, in order to improve the BH property in addition to the shape freezing property, by suppressing the addition amount of Al as much as possible, precipitation of AlN was suppressed and the amount of solid solution N was increased. As a result, it was found that an excellent BH property can be obtained without deteriorating the shape freezing property. In addition, the present inventors further studied production conditions for leaving unrecrystallized ferrite, and suppressed recrystallization by increasing the rate of temperature rise from the recrystallization temperature to the Ac ₁ transformation temperature in the annealing process. Further, the temperature of the steel sheet is changed to the Ac ₁ transformation so that the transformation to the austenite does not proceed excessively when heated to the α + γ two-phase region where the ferrite and austenite coexist, that is, the Ac ₁ transformation temperature or higher. It has been found that it is important to optimize the residence time that is above the temperature and the highest temperature reached for annealing.

The present invention has been made based on such findings, and the gist thereof is as follows.
(1) By mass%, C: 0.05 to 0.25%, Mn: 0.50 to 3.50%, Si: 1.00% or less, P: 0.150% or less, S: It is limited to 0.0150% or less, Al: 0.200% or less, N: 0.0100% or less, and Ti: 0.005 to 0.100%, Nb: 0.005 to 0.100% Or containing both, the balance is made of iron and inevitable impurities, the metal structure is made of ferrite and a hard second phase, the ferrite is made of recrystallized ferrite, one or both of transformation ferrite and unrecrystallized ferrite, The area ratio of non-recrystallized ferrite is 10 to 70%, the area ratio of one or both of the recrystallized ferrite and the transformed ferrite is 10 to 70%, and the area ratio of the hard second phase is 1 to 30%. %, And {5 in thickness 1/2 layer 54} <225>, {111} <112>, and {111} <110>, the average value of the X-ray random intensity ratios of the three crystal orientations is 3.5 or less. .

(2) By mass%, Mo: 0.1 to 1.5%, B: 0.0005 to 0.0100%, Cr: 0.10 to 1.50%, Ni: 0.10 to 1.50% Among them, the high-strength cold-rolled steel sheet according to (1) above, which contains one or two kinds.
(3) The high-strength cooling according to (1) or (2) above, characterized in that Al is limited to 0.0200% or less and N: 0.0010 to 0.0100% in mass%. Rolled steel sheet.
(4) The high-strength cold-rolled steel sheet according to (3), wherein the Al content and the N content satisfy Al / N ≦ 2.

(5) A high-strength cold-rolled steel sheet characterized by providing hot-dip Zn plating on the surface of the cold-rolled steel sheet according to any one of (1) to (4).
(6) A high-strength cold-rolled steel sheet characterized by providing alloyed hot-dip Zn plating on the surface of the cold-rolled steel sheet according to any one of (1) to (4).

(7) A steel slab having the chemical component according to any one of (1) to (4) above is hot-rolled, wound up in a temperature range of 300 to 500 ° C., pickled, and 60% or less. Cold rolling is performed at a reduction rate, and the steel sheet is heated at a rate of temperature increase from (Ac ₁ [° C.]-100 ° C.) to Ac ₁ [° C.] of 10 ° C./s or more, from Ac ₁ [° C.] to {Ac ₁ [ [° C.] + 2/3 × (Ac ₃ [° C.] − Ac ₁ [° C.])}, and annealing is performed with the residence time within the temperature range of the steel sheet being 10 to 200 s. A method for producing a high-strength cold-rolled steel sheet.
Here, Ac ₁ [° C.] and Ac ₃ [° C.] are the contents of C, Mn, and Si expressed in mass% (% C), (% Mn), and (% Si) according to the following (formula 1 ) and (a Ac ₁ transformation temperature was determined from equation 2) and Ac ₃ transformation temperature.
_{Ac 1 = 761.3 + 212 (%} C) -45.8 (% Mn) +16.7 (% Si)
... (Formula 1)
_{Ac 3 = 915-325.9 (% C)} -35.9 (% Mn) +31.4 (% Si)
... (Formula 2)

(8) A method for producing a high-strength cold-rolled steel sheet, which is cooled to 350 to 500 ° C. and then subjected to hot-dip Zn plating after the annealing described in (7).
(9) A method for producing a high-strength cold-rolled steel sheet, comprising performing heat treatment for 10 seconds or more in a temperature range of 450 to 600 ° C. after performing the hot-dip Zn plating according to (8).
(10) A high-strength cold-rolled steel sheet characterized by subjecting the cold-rolled steel sheet produced by the method according to any one of (7) to (9) above to 0.1 to 5.0% skin pass rolling. Manufacturing method.

  According to the present invention, it is possible to provide a cold-rolled steel sheet having a high tensile strength, which has excellent shape freezing property, preferably has a low yield strength, and more preferably has a paint bake hardenability, and a method for producing the same.

Conventionally, there has been no idea that a part of ferrite in the metal structure of a cold-rolled steel sheet is left as non-recrystallized ferrite and actively used for high strength. This is because if the recrystallization is incomplete, the material of the cold-rolled steel sheet is considered to be non-uniform, and the residual non-recrystallized ferrite is suppressed as much as possible.
Therefore, the conventional cold-rolled steel sheet composed of non-recrystallized ferrite and a hard second phase is composed of ferrite recrystallized during annealing heating (called recrystallized ferrite) in addition to non-recrystallized ferrite and austenite during cooling after annealing. Transformed ferrite (called transformed ferrite) is not mixed, and it is considered that the ferrite is only homogeneous unrecrystallized ferrite.

  Conventionally, a manufacturing method has been proposed in which the temperature rise rate of annealing is increased and the crystal grain size of the steel sheet is reduced, but this method completely eliminates unrecrystallized ferrite by holding in the α + γ two-phase region. It is thought that it was transformed into austenite. That is, in this prior art, after non-recrystallized ferrite is completely transformed to austenite by annealing, DP steel composed of ferrite re-transformed from austenite and a hard second phase during cooling remains unrecrystallized ferrite. It is estimated that it can be obtained without.

  However, when austenite is transformed into ferrite during cooling after annealing, austenite decomposes into ferrite and cementite. Therefore, a DP steel made of ferrite containing hard second phase made of bainite, martensite and retained austenite and cementite is obtained. Therefore, it is considered that the conventional DP steel obtained by increasing the rate of temperature increase during annealing is further promoted by cementite in reducing the local ductility.

  On the other hand, when unrecrystallized ferrite is actively left as in the present invention schematically shown in FIG. 1, a soft ferrite, that is, between recrystallized ferrite and transformed ferrite and the hard second phase, Unrecrystallized ferrite having the following strength can be present. The presence of non-recrystallized ferrite having intermediate strength between the soft ferrite and the hard second phase alleviates the concentration of strain at the interface between the ferrite and the hard second phase.

  Therefore, in the cold-rolled steel sheet of the present invention in which non-recrystallized ferrite is actively used, generation of voids generated at the interface between the soft ferrite and the hard second phase is suppressed. Furthermore, when non-recrystallized ferrite is actively left to suppress the formation of transformation ferrite, the generation of cementite that is the starting point of voids is also suppressed. Therefore, local ductility is remarkably improved, stretch flange formability is improved, and severe burring is possible.

  Non-recrystallized ferrite is one in which the crystal grains of ferrite stretched in the rolling direction by cold rolling are not recrystallized, and dislocations in the grains are recovered. Therefore, as schematically shown in FIG. 2, the grains of unrecrystallized ferrite often have subgrains formed by dislocation recovery. Further, in the grains of unrecombined ferrite, the crystal orientation continuously changes due to plastic deformation by cold rolling. On the other hand, in the recrystallized ferrite and the transformed ferrite, the crystal orientation in the grains becomes almost uniform by recrystallization or transformation, and the crystal orientations of adjacent crystal grains are greatly different.

In addition, the inventors have studied a method for leaving unrecrystallized ferrite,
(X) When the recrystallization temperature of the ferrite is lower than the Ac ₁ transformation temperature (hereinafter also referred to as Ac ₁ ), which is the temperature at which transformation from ferrite to austenite (referred to as α-γ transformation) starts, Increasing the rate of temperature rise from the recrystallization temperature to Ac ₁ ;
(Y) When the recrystallization temperature of ferrite is higher than Ac ₁ , recrystallization does not proceed regardless of the rate of temperature increase,
(Z) When the upper limit of the annealing temperature is too high, or when the residence time at Ac ₁ or higher is too long, the α-γ transformation proceeds and no unrecrystallized ferrite remains,
I found.

In order to leave unrecrystallized ferrite, annealing conditions are extremely important in the present invention, and it is particularly necessary to limit the rate of temperature rise below Ac ₁ , the maximum temperature reached, and the holding time above Ac _1. There is.

The temperature increase rate from (Ac ₁ [° C.] − 100 ° C.) to Ac ₁ [° C.] in annealing is 10 ° C./s or more. The lower limit of the temperature at which the rate of temperature increase is 10 ° C./s or higher is set to (Ac ₁ [° C.] − 100 ° C.) or higher. The lower limit of the recrystallization temperature of the DP steel of the present invention is increased by the content of components. This is because it is at least (Ac ₁ [° C.] − 100 ° C.). Further, the upper limit of the temperature of the heating rate and 10 ° C. / s or more was Ac ₁ [° C.] is the Ac ₁ [° C.] or higher temperatures occurs the alpha-gamma transformation, recrystallization is substantially stopped Because.

On the other hand, when the rate of temperature rise is less than 10 ° C./s, the area ratio of non-recrystallized ferrite is remarkably reduced by sufficiently proceeding recrystallization. In order to suppress the coarsening of the recrystallized ferrite, it is preferable to set the rate of temperature rise to more than 20 ° C./s. Furthermore, since steel with a low content of component has a low Ac ₁ , recrystallization proceeds more easily. When manufacturing such a steel with low Ac _1, it is preferable to set the rate of temperature rise to more than 30 ° C./s in order to ensure non-recrystallized ferrite.

Further, the lower limit of the maximum temperature achieved in annealing is set to Ac ₁ [° C.] or more, and the upper limit is set to {Ac ₁ [° C.] + 2/3 × (Ac ₃ [° C.] − Ac ₁ [° C.])}. When the maximum temperature reached is less than Ac ₁ , the ferrite does not transform to austenite, so the amount of the hard second phase is insufficient, and the strength-ductility balance is impaired. On the other hand, when the maximum temperature reaches {Ac ₁ [° C.] + 2/3 × (Ac ₃ [° C.] − Ac ₁ [° C.])}, the austenite transformation proceeds too much, so it is difficult to secure unrecrystallized ferrite. become.

Also, the residence time in the temperature range the temperature of the steel sheet is Ac ₁ [° C.] or higher and 10～200S. This is due to the following reason. That is, if the time for the temperature of the steel sheet to be Ac ₁ [° C.] or more is less than 10 s, the α-γ transformation does not proceed sufficiently, so that the hard second phase cannot be secured and the strength-ductility balance is impaired. On the other hand, if the residence time at Ac ₁ [° C.] or more exceeds 200 s, the austenite transformation proceeds too much, so that it is difficult to secure unrecrystallized ferrite.

Ac ₁ [° C.] and Ac ₃ [° C.] are the Ac ₁ transformation point and Ac ₃ transformation point, respectively, and are the contents of C, Mn, and Si expressed in mass% (% C), ( % Mn) and (% Si) are temperatures obtained from the following (formula 1) and (formula 2).

Ac ₁ = 761.3 + 212 (% C) −45.8 (% Mn) +16.7 (% Si) (Formula 1)
_{Ac 3 = 915-325.9 (% C)} -35.9 (% Mn) +31.4 (% Si) ··· ( Equation 2)

  Hereinafter, the reasons for limitation of the present invention will be described sequentially.

First, steel components will be described. In addition,% means the mass%.
Ti and Nb are the most important elements in the present invention, and one or both of them are contained. When one or both of Ti and Nb are added, a texture advantageous to shape freezing property is easily developed in the steel sheet after hot rolling, and recrystallization of processed ferrite is suppressed in the annealing process after cold rolling. Unrecrystallized ferrite can be left. Thereby, shape freezing property improves and high intensity | strength is obtained. In order to obtain such an effect, the lower limits of the Ti content and the Nb content must both be 0.005% or more. On the other hand, the upper limit of the Ti content and the Nb content is required to be 0.100% or less because the ductility is lowered when it exceeds 0.100%. From the viewpoint of alloy cost, the preferable upper limit of the Ti amount and the Nb amount is 0.040%.

  C is an element that promotes the formation of the hard second phase and contributes to an increase in strength, and an appropriate amount is added according to the target strength level. If the amount of C is less than 0.05%, it is difficult to obtain high strength, so the lower limit is made 0.05%. On the other hand, if the amount of C exceeds 0.25%, deterioration of formability and weldability is caused, so 0.25% is made the upper limit.

Si is a deoxidizing element, and the lower limit of the amount of Si is not specified, but if it is less than 0.01%, the manufacturing cost increases, so the lower limit is preferably made 0.01%. In addition, Si serves to increase the strength as a solid solution strengthening element and is also effective for obtaining a hard second phase. However, if the amount of Si exceeds 1.00%, Ac ₁ becomes too high, and it is necessary to increase the annealing temperature, and the transformation is promoted to make it difficult to secure unrecrystallized ferrite. % Or less. Further, if Si is added in excess of 0.50%, there may be a problem that the plating adhesion is deteriorated when performing hot dip Zn plating and the productivity is lowered due to the delay of the alloying reaction. Therefore, it is preferable that the upper limit of the Si amount is 0.50% or less.

Mn is an element that lowers Ac ₁ and the Ac ₃ transformation temperature (hereinafter also referred to as Ac ₃ ), which is the temperature at which the α-γ transformation is completed and becomes an austenite single phase, and is extremely important in the present invention. That is, when the amount of Mn is small, it is necessary to increase the annealing temperature, the transformation is promoted, and it becomes difficult to secure unrecrystallized ferrite. Mn, like Si, has an effect of increasing strength as an element contributing to solid solution strengthening, and is also effective for obtaining a hard second phase. From these viewpoints, the lower limit of the amount of Mn is 0.50%. On the other hand, if the amount of Mn exceeds 3.50%, the formability and weldability are deteriorated, so 3.50% is made the upper limit.

  P is an impurity and segregates at the grain boundary, which causes a reduction in toughness and weldability of the steel sheet. Furthermore, the alloying reaction is extremely slow during hot-dip Zn plating, and productivity is reduced. From these viewpoints, the upper limit of the P content is 0.150%. The lower limit is not particularly limited, but P is an element that enhances the strength at a low cost, so the P content is preferably 0.005% or more.

  S is an impurity, and if its content exceeds 0.0150%, hot cracking is induced or workability is deteriorated, so the upper limit is made 0.0150%.

  Al is a deoxidizer and does not define a lower limit, but is an element that remarkably increases the transformation point, so the upper limit is made 0.200%. From the viewpoint of manufacturing cost, it is preferable that the lower limit of the amount of Al is 0.0005% or more. Further, in order to obtain a sufficient deoxidation effect, it is preferable to add Al by 0.010% or more. Furthermore, in order to ensure the amount of dissolved N and to have BH properties, the upper limit of the Al amount is preferably 0.0200%. This is because Al is an element that forms AlN to reduce the amount of dissolved N.

  N is an impurity, and when the amount of N exceeds 0.0100%, the deterioration of toughness and ductility and the occurrence of cracks in the steel slab become remarkable. Note that N is effective for obtaining the hard second phase, and therefore may be positively added with an upper limit of 0.0100%. Furthermore, in order to secure the amount of dissolved N and to have BH properties, it is preferable to add N at 0.0010% or more.

  Furthermore, when Al is contained excessively with respect to the N content, the generation of AlN is promoted, and it may be difficult to secure the solid solution N amount. Therefore, in order to improve the BH property, the ratio Al / N between the Al content and the N content is preferably 2 or less. As described above, in order to improve the BH property, N is preferably added in an amount of 0.0010% or more, and it is desirable not to contain Al, that is, to make the content 0%. Therefore, it is desirable that the lower limit of Al / N be 0. However, from the viewpoint of manufacturing cost, the preferable lower limit of Al is 0.0005% as described above. Moreover, since the upper limit of N amount is 0.0100%, 0.05 is optimal as the lower limit of Al / N.

In addition, when Ac ₁ reaches a high temperature of 700 ° C. or higher, the α-γ transformation proceeds in a short time during annealing in the α + γ two-phase region. Therefore, in the present invention, Ac ₁ is 700 ° C. or lower. It is preferable that

  Since all of Mo, B, Cr and Ni are elements that enhance the hardenability, one or more of them may be added as necessary. In order to obtain high strength by improving hardenability, Mo: 0.1% or more, B: 0.0005% or more, Cr: 0.1% or more, Ni: 0.1% or more are added. preferable. On the other hand, excessive addition leads to an increase in alloy cost, so it is preferable that Mo: 1.5% or less, B: 0.01% or less, Cr: 1.5% or less, Ni: 1.5% or less. .

Next, the microstructure and texture will be described.
The microstructure of the steel sheet obtained by the present invention is composed of ferrite and a hard second phase, and the ferrite is a general term for non-recrystallized ferrite, recrystallized ferrite and transformed ferrite. In addition, in the structure observation with an optical microscope, the difference between recrystallized ferrite and transformed ferrite is not clear, and it is difficult to distinguish them.

  The hard second phase consists of martensite, bainite and pearlite and may contain less than 3% retained austenite. While the hard second phase contributes to high strength, if it is present in excess, the ductility is remarkably lowered, so the lower limit is 1% and the upper limit is 30%.

  The microstructure may be obtained by taking a sample with the cross section of the plate thickness parallel to the rolling direction as the observation surface, polishing the observation surface, performing nital etching, and if necessary, repeller etching, and observing with an optical microscope. By analyzing the microstructure image obtained by the optical microscope, the total amount of one or more of pearlite, bainite and martensite is obtained as the area ratio of the phase other than ferrite. be able to. Although it is difficult to distinguish residual austenite from martensite with an optical microscope, the volume ratio can be measured by an X-ray diffraction method. Note that the area ratio obtained from the microstructure is the same as the volume ratio.

  The area ratio of one or both of the recrystallized ferrite and the transformed ferrite is 10 to 70%. This is because the ductility decreases when the area ratio of one or both of the recrystallized ferrite and the transformed ferrite is less than 10%, and the strength decreases when the area ratio exceeds 70%.

  Since non-recrystallized ferrite contributes to high strength, it is necessary to contain 10% or more of non-recrystallized ferrite in order to obtain the effect. On the other hand, if the area ratio of non-recrystallized ferrite exceeds 70%, the ductility is remarkably lowered, so the upper limit is made 70%.

  Non-recrystallized ferrite and other ferrites, that is, recrystallized ferrite and transformed ferrite, are obtained by analyzing the crystal orientation measurement data of an electron back scattering pattern (EBSP) by the Kernel Average Misoration method (KAM method). It can be determined by analysis.

  In the grains of unrecrystallized ferrite, although dislocations are recovered, there is a continuous change in crystal orientation caused by plastic deformation during cold rolling. On the other hand, the crystal orientation change in the ferrite grains excluding non-recrystallized ferrite becomes extremely small. This is because the crystal orientation does not change in one crystal grain, although the crystal orientation of adjacent crystal grains varies greatly due to recrystallization and transformation. In the KAM method, the crystal orientation difference between adjacent pixels (measurement points) can be quantitatively shown. Therefore, in the present invention, the average crystal orientation difference between adjacent measurement points is within 1 ° and the average crystal orientation difference is 2 When a pixel boundary is defined as a grain boundary, a grain having a crystal grain size of 3 μm or more is defined as ferrite other than unrecrystallized ferrite, that is, recrystallized ferrite and transformed ferrite.

  The EBSP measurement may be performed in a range of 100 × 100 μm at a 1/4 thickness position in the plate thickness direction of any plate cross section at a measurement interval of 1/10 of the average crystal grain size of the sample after annealing. As a result of the EBSP measurement, the measurement points obtained are output as pixels. Samples to be used for EBSP crystal orientation measurement are prepared so that the steel plate is reduced to a predetermined thickness by mechanical polishing, etc., and then the strain is removed by electrolytic polishing, etc., and at the same time, the 1/4 thickness is the measurement surface. To do.

  Since the total area ratio of ferrite including non-recrystallized ferrite is the remainder of the area ratio of the hard second phase, the sample used for measuring the crystal orientation of EBSP was nital etched, and the optical microscope of the field of view where the measurement was performed The photograph may be taken at the same magnification, and the obtained tissue photograph may be obtained by image analysis. Furthermore, by comparing the result of the crystal orientation measurement of this structural photograph and EBSP, the total area ratio of non-recrystallized ferrite and ferrite other than non-recrystallized ferrite, that is, recrystallized ferrite and transformed ferrite can be obtained. .

  The average value of the X-ray random intensity ratios of the three crystal orientations of {554} <225>, {111} <112> and {111} <110> in the 1/2 layer thickness of the steel of the present invention is 3.5 or less. It is preferable that As this orientation develops, the shape freezeability deteriorates, so the upper limit is set to 3.5.

  Note that the X-ray random intensity ratio means that the X-ray intensity of a standard sample that does not accumulate in a specific orientation and the test material is measured under the same conditions by the X-ray diffraction method or the like. It is a numerical value obtained by dividing the line intensity by the X-ray intensity of the standard sample.

  The X-ray random intensity ratio in the {554} <225>, {111} <112> and {111} <110> orientations is measured by X-ray diffraction {110}, {100}, {211}, {310 } Of the pole figures, a three-dimensional texture (ODF) calculated by a series expansion method using a plurality of pole figures may be used. That is, in order to obtain the X-ray random intensity ratio of the {554} <225>, {111} <112> and {111} <110> orientations, (554) [1-10] in the φ2 = 45 ° cross section of the ODF And the intensity of (111) [1-10].

Samples to be subjected to X-ray diffraction are prepared so that the steel plate is reduced to a predetermined thickness by mechanical polishing, etc., and then distortion is removed by chemical polishing, electrolytic polishing, etc., and at the same time, the 1/2 surface thickness becomes the measurement surface To do. When there is a segregation zone or a defect in the thickness center layer of the steel plate, causing inconvenience in measurement, the above method is used so that an appropriate surface becomes the measurement surface in the range of 3/8 to 5/8 of the plate thickness. The sample may be adjusted according to the above and measured. If more difficult X-ray diffraction performs a statistical measure sufficient number of the EBSP method or ECP (E lectron C hanneling P attern ) method.

  Here, {hkl} <uvw> means that the normal direction of the plate surface is parallel to {hkl} and the rolling direction is parallel to <uvw> when an X-ray sample is collected by the above-described method. Is shown. The crystal orientation is usually indicated by [hkl] or {hkl}, which is perpendicular to the plate surface, and (uvw) or <uvw>, which is parallel to the rolling direction. {Hkl} and <uvw> are generic terms for equivalent planes, and [hkl] and (uvw) indicate individual crystal planes. That is, in the present invention, b. c. c. Since the structure is targeted, for example, (111), (−111), (1-11), (11-1), (−1-11), (-11-1), (1-1-1) , (-1-1-1) planes are equivalent and indistinguishable. In such a case, these orientations are collectively referred to as {111}. Since the ODF display is also used to display the orientation of other crystal structures with low symmetry, the individual orientation is generally displayed as [hkl] (uvw). However, in the text, [hkl] (uvw) ) And {hkl} <uvw> are synonymous.

Next, the reason for limiting the manufacturing conditions will be described.
The slab used for hot rolling is not particularly limited. That is, what was manufactured with the continuous casting slab, the thin slab caster, etc. should just be used. It is also compatible with processes such as continuous casting-direct rolling, in which hot rolling is performed immediately after casting.

  Conditions such as the rolling temperature and rolling reduction in hot rolling need not be specified, and may be performed according to a conventional method. Cooling up to winding after hot rolling may be performed as appropriate, such as forced cooling with water, refrigerant, blowing, mist, or the like.

  The winding temperature after finish rolling is important in the present invention, and is set to 300 to 500 ° C. Thereby, the structure of the steel sheet after hot rolling becomes a structure in which hard phases such as martensite and pearlite are dispersed in soft polygonal ferrite, and {554} <225>, {111} <112 during cold rolling. > And {111} <110> orientation can be prevented from developing. When the coiling temperature exceeds 500 ° C., a hard phase such as martensite or pearlite cannot be obtained, and the texture after cold rolling annealing deteriorates. Therefore, the upper limit of the coiling temperature is set to 500 ° C. On the other hand, if the winding is performed at a temperature lower than 300 ° C., the precipitation amount of Ti carbonitride in the hot-rolled sheet is reduced, so that the recrystallization suppressing effect is reduced. Therefore, the lower limit of the winding temperature is set to 300 ° C.

  The hot-rolled steel sheet thus manufactured is pickled and then cold-rolled at a rolling reduction of 60% or less. This is because if the cold rolling rate exceeds 60%, the {554} <225>, {111} <112>, and {111} <110> orientations that are disadvantageous to shape freezing properties develop.

Annealing is the most important step in the present invention, and the conditions are as described above. Annealing is preferably performed by continuous annealing equipment in order to control the rate of temperature rise and the heating time. Further, in order to increase the rate of temperature rise, a high-frequency heating device or an electric heating device may be used in combination. In annealing, the residence time in the Ac ₁ or more, the sum of the time the temperature of the steel sheet is Ac ₁ or more, the length of the set temperature and the furnace of the heating furnace can be controlled by the sheet passing speed.

  Further, the cooling rate after annealing is not particularly specified, but if the cooling rate is less than 1 ° C./s, a sufficiently hard second phase may not be obtained. From this viewpoint, the lower limit of the cooling rate is preferably 1 ° C./s. On the other hand, in order to make the cooling rate over 250 ° C./s, it is necessary to introduce special equipment and the like, so it is preferable to set the upper limit of the cooling rate to 250 ° C./s. The cooling rate after annealing may be appropriately controlled by forced cooling with water or the like, blowing of refrigerant, blowing air, mist, or the like.

  After annealing, an overaging treatment, hot dip Zn plating, or alloyed hot dip Zn plating may be performed as necessary. The composition of the Zn plating is not particularly limited, and in addition to Zn, Fe, Al, Mn, Cr, Mg, Pb, Sn, Ni, or the like may be added as necessary. The plating may be performed in a separate process from the annealing, but from the viewpoint of productivity, it is preferable that the plating is performed by a continuous annealing-hot Zn plating line in which annealing and plating are continuously performed. Also in this case, in order to ensure non-recrystallized ferrite, it is necessary to perform annealing under the above conditions.

  When performing an alloying process, it is preferable to carry out in the temperature range of 450-600 degreeC. If it is less than 450 ° C, alloying does not proceed sufficiently, and if it exceeds 600 ° C, alloying proceeds excessively and the plating layer becomes brittle, which causes problems such as peeling of the plating by processing such as pressing. . When the alloying treatment time is less than 10 s, alloying may not proceed sufficiently. Further, the upper limit of the alloying time is not particularly defined, but is preferably within 100 s from the viewpoint of production efficiency.

  Further, from the viewpoint of productivity, it is preferable to continuously provide an alloying treatment furnace in the continuous annealing-hot Zn plating line and to perform annealing, plating, and alloying treatment continuously.

Steel having the composition shown in Table 1 was melted in a vacuum melting furnace, and hot rolling, cold rolling and annealing were performed under the conditions shown in Table 2. The average cooling rate from the soaking temperature to 500 ° C. or the overaging temperature was 50 ° C./s. Here, [-] in Table 1 means that the analysis value of the component was less than the detection limit. Table 1 also shows the calculated values of Ac ₁ [° C.] and Ac ₃ [° C.].

In Table 2, CT [° C.] is a coiling temperature in the hot rolling process. The heating rate in Table 2 was calculated according to the time required for the temperature increase from (Ac ₁ [° C.]-100 ° C.) to Ac ₁ [° C.]. The residence time in Table 2 is the time during which annealing was performed in a temperature range of Ac ₁ [° C.] or higher.

  Among the cold-rolled steel sheets shown in Table 2, production No. 2 and 6 were subjected to an alloying treatment at 500 ° C. for 20 s after immersion in a Zn plating bath after the annealing step. Furthermore, among the cold-rolled steel sheets shown in Table 2, the production No. No. 11 was cooled from a soaking temperature to 300 ° C. at a cooling rate of 50 ° C./s as described above, held at that temperature for 400 s, then cooled to room temperature at 10 ° C./s, and mechanical properties were measured.

  From the annealed cold-rolled steel sheet, a No. 5 tensile test piece of JIS Z 2201 was taken with the width direction (TD direction) perpendicular to the rolling direction as the longitudinal direction, and in accordance with JIS Z 2241, tensile properties in the TD direction Evaluated. Further, the X-ray random intensity ratio of {554} <225>, {111} <112> and {111} <110> orientations was measured by X-ray diffraction, and an average value of three crystal orientations was obtained. Hereinafter, the average value of the X-ray random intensity ratios of crystal orientations of {554} <225>, {111} <112>, and {111} <110> is referred to as an average random intensity ratio. A sample for X-ray diffraction was prepared by mechanical polishing and electrolytic polishing so that the half-thickness surface was the measurement surface.

  The cross-sectional microstructure was observed with an optical microscope, and the total area ratio of pearlite, bainite, and martensite was obtained by image analysis of a structural photograph taken with an optical microscope. Moreover, the area ratio of the non-recrystallized ferrite and the area ratio of the ferrite excluding the non-recrystallized ferrite were obtained by collating the crystal orientation measurement of EBSP, the measurement result, and the optical microscope structure photograph, and image analysis.

  The evaluation of the shape freezing property was made from a cold-rolled steel sheet after annealing by preparing a test piece having a width of 25 mm and a length of 130 mm and using a die having a punch width of 40 mm, a punch shoulder R5, a die width of 43.4 mm and a die shoulder R3 After forming into a hat shape having a height of 40 mm with various wrinkle holding pressures, the amount of warpage of the wall was measured as a local radius ρ (mm), and the reciprocal was 1000 / ρ. The smaller the 1000 / ρ, the better the shape freezing property.

  Table 3 shows the evaluation of the area ratio of the microstructure, the average X-ray random intensity ratio, the mechanical properties, and the shape freezing property. In the column of “Evaluation of shape freezing property”, the evaluation result is expressed as “O” when 1000 / ρ ≦ 0.017 × TS (MPa) −4.13, and “X” in other cases. Described.

  The steel having the chemical composition of the present invention is hot-rolled and cold-rolled under appropriate conditions, and after an appropriate annealing process, it has a high strength with a low yield ratio even if it is over-aged, Zn plated, and further alloyed. It is possible to obtain a cold-rolled steel sheet.

  On the other hand, Steel No. Since B has a small amount of C, its strength is lowered. Steel No. D has a large amount of Si. F has less Mn and needs to be annealed at a high temperature. Since J has less Ti, non-recrystallized ferrite is reduced, the average random strength ratio is increased, and the yield ratio is decreased.

  In addition, production No. No. 3 has a high winding temperature in the hot rolling process, and the production No. 3 No. 5 has a high cold rolling reduction ratio, a high average random strength ratio, and a high yield ratio.

Production No. _No. 10 has a slow rate of temperature increase from (Ac ₁ [° C.] − 100 ° C.) to Ac ₁ [° C.] and less unrecrystallized ferrite, resulting in a higher average random strength ratio and a higher yield ratio. . Production No. _No. 14 has a long residence time at Ac ₁ [° C.] or higher, so that there is little unrecrystallized ferrite and high strength, but the local elongation is lowered, the average random strength ratio is increased, and the yield ratio is increased. Yes.

  Production No. No. 15 has a low maximum temperature for annealing and a hard second phase could not be obtained, so the strength was reduced and the yield ratio was large. Production No. No. 18 has a high maximum temperature for annealing, less unrecrystallized ferrite, increased hard second phase, high strength, but reduced local elongation, increased average random strength ratio, and increased yield ratio It has become.

Steel having the composition shown in Table 4 was melted in a vacuum melting furnace, and hot rolling, cold rolling and annealing were performed under the conditions shown in Table 5. The average cooling rate from the soaking temperature to 500 ° C or the overaging temperature was 50 ° C / s. Here, [-] in Table 4 means that the analysis value of the component was less than the detection limit. Table 4 also shows the calculated values of Ac ₁ [° C.] and Ac ₃ [° C.].

In Table 5, CT [° C.] is the coiling temperature in the hot rolling process. The heating rate in Table 5 was calculated according to the time required for the temperature increase from (Ac ₁ [° C.]-100 ° C.) to Ac ₁ [° C.]. The residence time in Table 5 is the time during which annealing was performed in a temperature range of Ac ₁ [° C.] or higher.

  Among the cold-rolled steel sheets shown in Table 5, the production No. About 21 and 24, the alloying process was performed for 20 s at 500 degreeC after immersion in Zn plating bath after the annealing process. Furthermore, among the cold-rolled steel sheets shown in Table 5, the production No. As for No. 28, it was cooled from the soaking temperature to 300 ° C. at the cooling rate of 50 ° C./s as described above, held at that temperature for 400 s, then cooled to room temperature at 10 ° C./s, and the mechanical properties were measured.

  After annealing, the tensile properties and shape freezing properties in the TD direction were evaluated in the same manner as in Example 1, and X-ray randomization of crystal orientations of {554} <225>, {111} <112>, {111} <110> The intensity ratio was measured to determine the average random intensity ratio. Further, the total area ratio of pearlite, bainite and martensite, the area ratio of unrecrystallized ferrite, and the area ratio of ferrite excluding non-recrystallized ferrite were also determined in the same manner as in Example 1. Furthermore, the amount of BH was evaluated based on the paint bake hardening test described in the annex of JIS G 3135.

  Table 6 shows the evaluation of the area ratio of the microstructure, the average X-ray random intensity ratio, the mechanical properties, the BH property, and the shape freezing property.

  The steel having the chemical components of the present invention is hot-rolled and cold-rolled under appropriate conditions and subjected to an appropriate annealing step, so that even if it is over-aged, Zn plated, and further alloyed, it has a high low yield ratio. It is possible to obtain a strength cold-rolled steel sheet. Furthermore, the bake hardenability is also improved by reducing the Al amount and Al / N. Of the examples of the present invention, the production No. No. 33 is a steel no. This is an example of AF. Nos. 35 and 36 are steel Nos. With an Al content exceeding 0.020%. It is an example of AG. In these, the area ratio of the microstructure and the average X-ray random intensity ratio are within the scope of the present invention, and the mechanical properties are excellent, but the BH amount is slightly decreased.

On the other hand, production No. No. 22 has a high winding temperature in the hot rolling step, and the production No. 22 No. 34 has a high cold rolling reduction ratio, a high average random strength ratio, and a high yield ratio. Production No. In _No. 24, the rate of temperature increase from (Ac ₁ [° C.] − 100 ° C.) to Ac ₁ [° C.] was slow and the amount of unrecrystallized ferrite decreased, so the average random strength ratio increased and the yield ratio increased. . Production No. _No. 29 has a long residence time at Ac ₁ [° C.] or higher, so that there is little unrecrystallized ferrite and high strength, but the local elongation is lowered, the average random strength ratio is increased, and the yield ratio is increased. Yes.

  Production No. No. 30 has a low maximum temperature for annealing and a hard second phase could not be obtained, so the strength was lowered and the yield ratio was increased. Production No. No. 26 has a high maximum temperature for annealing, less unrecrystallized ferrite, increased hard second phase, and high strength, but reduced local elongation, increased average random strength ratio, and increased yield ratio It has become.

  In addition, production No. No. 37 is a steel No. having a small amount of C and Mn. This is an example of AH, and the strength is reduced. In addition, production No. No. 38 is a steel no. This is an example of AI, and it is necessary to perform annealing at a high temperature. No. 39 is a steel no. These are examples of AJ, each of which has less unrecrystallized ferrite, an average random strength ratio is increased, and a yield ratio is decreased. Furthermore, steel no. Since AH, AI, and AJ have a large amount of Al, production No. In Nos. 37 to 39, the solute N amount cannot be sufficiently secured, and the BH property is also deteriorated.

It is a schematic diagram of the metal structure of the steel of this invention. It is a schematic diagram of the non-recrystallized ferrite of the present invention.

Explanation of symbols

1 Unrecrystallized ferrite 2 Hard second phase 3 Recrystallized ferrite or transformation ferrite 4 Subgrain

Claims (10)

% By mass
C: 0.05 to 0.25%,
Mn: 0.50 to 3.50%
Containing
Si: 1.00% or less,
P: 0.150% or less,
S: 0.0150% or less,
Al: 0.200% or less,
N: limited to 0.0100% or less, and
Ti: 0.005 to 0.100%,
Nb: 0.005 to 0.100%
One or both of the above, the balance being composed of iron and inevitable impurities, the metallographic structure is composed of ferrite and a hard second phase, and the ferrite is composed of recrystallized ferrite, one or both of transformed ferrite and unrecrystallized ferrite The area ratio of the non-recrystallized ferrite is 10 to 70%, the area ratio of one or both of the recrystallized ferrite and the transformed ferrite is 10 to 70%, and the area ratio of the hard second phase is 1 The average value of the X-ray random intensity ratios of the three crystal orientations of {554} <225>, {111} <112> and {111} <110> in the ½ layer thickness is 30%. A high-strength cold-rolled steel sheet characterized by being 5 or less. % By mass
Mo: 0.1 to 1.5%,
B: 0.0005 to 0.0100%,
Cr: 0.10 to 1.50%,
Ni: 0.10 to 1.50%
Among them, the high-strength cold-rolled steel sheet according to claim 1, further comprising one or more of them. % By mass
Al: limited to 0.0200% or less,
N: 0.0010 to 0.0100%
The high-strength cold-rolled steel sheet according to claim 1 or 2, characterized by comprising: Al content and N content are
Al / N ≦ 2
The high-strength cold-rolled steel sheet according to claim 3, wherein:   A high-strength cold-rolled steel sheet, wherein hot-dip Zn plating is provided on the surface of the cold-rolled steel sheet according to any one of claims 1 to 4.   A high-strength cold-rolled steel sheet characterized by providing alloyed hot-dip Zn plating on the surface of the cold-rolled steel sheet according to any one of claims 1 to 4. A steel slab having the chemical composition according to any one of claims 1 to 4 is hot-rolled, wound in a temperature range of 300 to 500 ° C, pickled, and then cold-rolled at a reduction rate of 60% or less. And the obtained steel sheet was heated at a rate of temperature increase from (Ac ₁ [° C.] − 100 ° C.) to Ac ₁ [° C.] of 10 ° C./s or more, from Ac ₁ [° C.] to {Ac ₁ [° C.] + 2 / 3 × (Ac ₃ [° C.] − Ac ₁ [° C.])}, and annealing is performed with the residence time within which the temperature of the steel sheet is within the temperature range being 10 to 200 s. A method for producing a high-strength cold-rolled steel sheet.
Here, Ac ₁ [° C.] and Ac ₃ [° C.] are the contents of C, Mn, and Si expressed in mass% (% C), (% Mn), and (% Si) according to the following (formula 1 ) and (a Ac ₁ transformation temperature was determined from equation 2) and Ac ₃ transformation temperature.
_{Ac 1 = 761.3 + 212 (%} C) -45.8 (% Mn) +16.7 (% Si)
... (Formula 1)
_{Ac 3 = 915-325.9 (% C)} -35.9 (% Mn) +31.4 (% Si)
... (Formula 2)   A method for producing a high-strength cold-rolled steel sheet, characterized by cooling to 350 to 500 ° C after the annealing according to claim 7 and then applying hot-dip Zn plating.   A method for producing a high-strength cold-rolled steel sheet, comprising performing heat treatment for 10 seconds or more in a temperature range of 450 to 600 ° C after the hot-dip Zn plating according to claim 8 is performed.   A low yield ratio and high strength excellent in shape freezing property, characterized by subjecting a cold-rolled steel sheet produced by the method according to any one of claims 7 to 9 to 0.1 to 5.0% skin pass rolling. A method for producing a cold-rolled steel sheet.

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