JP5151468B2 - High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet suitable for applications such as automotive steel sheets and the like and excellent in workability such as strength-ductility balance and stretch flangeability, and excellent in impact resistance, and a method for producing the same.

  In order to reduce the amount of carbon dioxide emission, the weight reduction of automobile bodies aimed at 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 a high-strength steel sheet to an automobile body, excellent workability is also required. As a steel material having both such strength and workability, a dual-phase steel (Dual Phase steel, hereinafter referred to as DP steel) having a composite structure composed of a hard second phase mainly composed of ferrite and martensite is known. Yes.

  However, although conventional DP steel is superior in balance between strength and ductility (hereinafter also referred to as strength-ductility balance) compared to general-purpose steel, it is the product of tensile strength and total elongation (hereinafter TS). It was difficult to make x El) 18000 (MPa ·%) or more. Furthermore, conventional DP steel has low local elongation and poor stretch flangeability because microvoids are easily generated at the boundary between the soft phase ferrite and the hard phase martensite due to the hardness difference of each phase. There is a problem that the impact strength is inferior because the yield strength is low.

  In order to solve such a problem, steel sheets that have both strength-ductility balance and stretch flangeability have been proposed by refining the grain size of recrystallized ferrite after annealing (for example, Patent Documents 1 to 4). ).

  However, although Patent Document 1 describes that the local ductility is improved by finely dispersing the hard second phase uniformly, in the examples, the knowledge about the material properties of stretch flangeability such as local ductility is described. Is not shown.

  Further, the cold-rolled steel sheets proposed in Patent Documents 2 and 3 make the crystal grains of recrystallized ferrite very fine, the temperature range of recrystallization annealing after cold rolling is very narrow, Temperature control is extremely difficult.

  Furthermore, the cold-rolled steel sheet proposed in Patent Document 4 is forcibly cooled by spraying or blowing air of cooling water while immersing the coil after hot rolling in the cooling water or rewinding the coil. Is damaged.

  Patent Document 5 proposes a high-strength cold-rolled steel sheet composed of non-recrystallized ferrite and a hard second phase. Although the strength is high and the yield ratio is high, the elongation is low and the formability is insufficient. Met.

  In addition, some of the present inventors have proposed a high-strength steel sheet in combination with non-recrystallized ferrite in addition to solid solution strengthening and precipitation strengthening in order to strengthen the structure in Patent Document 6, It does not actively utilize unrecrystallized ferrite. Furthermore, in the patent documents 7 to 9, the present inventors have proposed a steel plate composed of a ferrite phase containing non-recrystallized ferrite and a hard second phase, and a method for producing the same. It is not intended to improve strength.

  On the other hand, the Young's modulus of the steel sheet has a correlation with the rigidity, and the Young's modulus is increased to secure the rigidity, and further, a steel sheet that has been strengthened by utilizing a hard second phase such as martensite and bainite has been proposed. (For example, Patent Documents 10 and 11). However, Patent Documents 10 and 11 do not show any knowledge about local ductility.

JP 2002-235145 A Japanese Patent Laid-Open No. 2003-27043 JP 2004-250774 A JP 2005-179732 A Japanese Patent Laid-Open No. 53-5018 JP 2006-283156 A Japanese Patent Application No. 2006-262873 Japanese Patent Application No. 2006-264253 Japanese Patent Application No. 2006-268647 JP 2005-314792 A JP 2005-314793 A

  The object of the present invention is to stably deteriorate the productivity of a high-strength cold-rolled steel sheet that has excellent impact resistance properties, more preferably high rigidity, in addition to workability such as strength-ductility balance and stretch flangeability. To provide without.

The present invention actively utilizes non-recrystallized ferrite, which is harder than recrystallized ferrite and softer than the hard second phase, and fine carbides of Nb and Ti, and has both collision resistance and local ductility. An improved high-strength steel sheet. When Nb and Ti are positively added, the temperature difference from the recrystallization temperature to the Ac ₁ transformation temperature is reduced, so that the temperature is close to the recrystallization temperature, that is, 100 ° C. from the Ac ₁ transformation temperature. This is based on the knowledge that it is important to appropriately set the rate of temperature increase from a low temperature to the Ac ₁ transformation temperature in order to appropriately retain unrecrystallized ferrite. More preferably, in the present invention, the hot rolled steel sheet having a texture developed in the hot rolling process is subjected to cold rolling exceeding 60%, and the {112} <110> orientation is developed to have a Young's modulus. It is an improved high strength steel plate. The gist of the present invention is as follows.

(1) By mass%, C: 0.05 to 0.25%, Si: 0.01 to 1.50%, Mn: 0.50 to 3.50%, P: 0.150% or less, S: 0.0150% or less, Al: 0.200% or less, N: 0.0100% or less, Nb: 0.005 to 0.100%, Ti: 0.005 to 0.100% Alternatively, both are contained in a total of less than 0.130%, the balance is made of iron and inevitable impurities, Ac ₁ [° C.] is 700 ° C. or more, the metal structure is made of ferrite and a hard second phase, It consists of non-recrystallized ferrite and one or both of recrystallized ferrite and transformed ferrite, the area ratio of the non-recrystallized ferrite is 10 to 70%, and the area ratio of one or both of the recrystallized ferrite and transformed ferrite is 10 to 70%, the hard second Wherein the area ratio of 1 to 30% local ductility, high-strength cold-rolled steel sheet excellent in workability and anti-collision properties.

(2) Further, in terms of mass%, Mo: 0.1 to 1.5%, B: 0.0005 to 0.0100%, Cr: 0.10 to 1.50%, Ni: 0.10 to 1. The high-strength cold-rolled steel sheet having excellent local ductility, workability, and impact resistance as described in (1) above, comprising 50% or more of one or more.

(3) The local ductility according to any one of (1) and (2) above, wherein the pole density in the {112} <110> orientation in the 1/2 layer thickness is 6 or more , A high-strength cold-rolled steel sheet with excellent workability and impact resistance.

(4) A high excellent in local ductility, workability and impact resistance, characterized in that hot-dip Zn plating is provided on the surface of the cold-rolled steel sheet according to any one of (1) to (3) above. Strength cold-rolled steel sheet.
(5) Excellent in local ductility, workability and impact resistance characteristics, characterized by providing alloyed hot-dip Zn plating on the surface of the cold rolled steel sheet according to any one of (1) to (3) above High strength cold rolled steel sheet.

(6) A steel slab having the chemical component according to any one of (1) or (2) above is hot-rolled, pickled and then cold-rolled, and then the steel sheet is (Ac ₁ [ [C]]-100 ° C.) to Ac ₁ [° C.] at a rate of 0.1-20 ° C./s, Ac ₁ [° C.]-{Ac ₁ [° C.] + 2/3 × (Ac ₃ [° C.] − the temperature was raised to Ac ₁ [℃]) within a temperature range of}, characterized in that the temperature of the steel sheet to annealing the residence time is within the temperature range as 10～200S, local ductility, workability and crashworthiness A method for producing high-strength cold-rolled steel sheets with excellent characteristics.
Here, Ac ₁ [° C.] and Ac ₃ [° C.] are expressed by mass% of C, Mn, and Si (% C), (% Mn), and (% Si) according to the following (formula 1) and These are the Ac ₁ transformation temperature and Ac ₃ transformation temperature determined from the 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)
... (Formula 2)

(7) A steel slab having the chemical component according to any one of (1) or (2) above, wherein the finish rolling temperature is not less than the Ar ₃ transformation temperature, and the reduction is performed in a range from 950 ° C. to the finish rolling temperature. The total ductility is 30% or more, hot rolling is performed, and after pickling, the steel sheet is annealed by cold rolling at a reduction rate of more than 60% , and the local ductility according to (6) above A method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance.

(8) After the annealing described in any one of (6) and (7) above, cooling to 350 to 500 ° C. and then applying hot-dip Zn plating , local ductility, workability and collision resistance A method for producing high-strength cold-rolled steel sheets with excellent characteristics.
(9) High strength excellent in local ductility, workability and impact resistance , characterized by performing heat treatment for 10 seconds or more in a temperature range of 450 to 600 ° C. after performing hot-dip Zn plating as described in (8) above A method for producing a cold-rolled steel sheet.

(10) The local ductility and processing , characterized by subjecting the cold-rolled steel sheet produced by the method according to any one of (6) to (9) above to 0.1 to 5.0% skin pass rolling. For producing a high-strength cold-rolled steel sheet excellent in heat resistance and impact resistance.

  According to the present invention, it becomes possible to provide a high-strength cold-rolled steel sheet having excellent workability and impact resistance, and more preferably excellent rigidity. In particular, non-recrystallized ferrite that can be stably manufactured without impairing productivity is positively provided. It is possible to provide a high-strength cold-rolled steel sheet that has excellent strength-ductility balance, stretch flangeability, impact resistance, and rigidity, and that contributes significantly to the industry.

  Conventionally, there has been no idea that a part of the ferrite of the metal structure of the cold-rolled steel sheet remains as non-recrystallized ferrite. This is because the material of the cold-rolled steel sheet is considered to be non-uniform when recrystallization is incomplete.

  Therefore, the conventional cold-rolled steel sheet composed of unrecrystallized ferrite and a hard second phase is composed of ferrite recrystallized during annealing (called recrystallized ferrite) and ferrite transformed from austenite during cooling after annealing (called transformed ferrite). )) And non-recrystallized ferrite are not mixed, and the ferrite is considered to be only homogeneous non-recrystallized ferrite.

  In addition, manufacturing methods have been proposed in which the temperature rise rate of annealing is increased and the crystal grain size of the steel sheet is made finer. However, these methods completely transform unrecrystallized ferrite into austenite by holding in the α + γ two-phase region. It is thought that it was what made it. 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, in the conventional DP steel obtained by increasing the heating rate during annealing, it is considered that the decrease in local ductility was further promoted by cementite.

  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 non-recrystallized 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 the present invention, precipitation strengthening by adding Nb and Ti in addition to non-recrystallized ferrite is utilized to increase yield strength and improve impact resistance. However, excessive addition of non-recrystallized ferrite has become a problem due to the addition of Nb and Ti, which are elements that suppress recrystallization. In order to prevent this, the optimum manufacturing conditions are studied, the components are adjusted so that the Ac ₁ transformation temperature is increased within the range that does not deteriorate the workability, and the temperature rise from the recrystallization temperature to the Ac ₁ transformation temperature in the annealing process. We found it important to slow down slightly.

Furthermore, the present inventors have noted that non-recrystallized ferrite has a cold-rolled texture formed by cold rolling as it is. That is, by optimizing the steel components, hot rolling and cold rolling, a cold-rolled steel sheet having a high Young's modulus can be obtained by developing a working texture that increases the Young's modulus and suppressing recrystallization due to annealing. We thought that we could do it and examined it. As a result, in order to obtain a cold-rolled steel sheet having an increased Young's modulus, the present inventors have a content of either one or both of Nb and Ti in the steel component composition, and for hot rolling, at the finishing temperature. It has been found that it is important that the rolling reduction in the temperature range in the vicinity of a certain Ar ₃ transformation temperature and that the cold rolling reduction be more than 60%. The Ar ₃ transformation temperature is a temperature at which ferrite transformation starts at the time of cooling, and is also simply referred to as Ar ₃ [° C.].

Further, in order to appropriately leave unrecrystallized ferrite, the α + γ two-phase region where the ferrite and austenite coexist, that is, the progress of transformation to austenite when heated to the Ac ₁ transformation temperature or higher. It is also important to suppress. Therefore, it was simultaneously found that it is important to optimize the residence time in which the temperature of the steel sheet is equal to or higher than the Ac ₁ transformation temperature and the highest temperature reached for annealing.

The heating rate from (Ac ₁ [° C.] − 100 ° C.) to Ac ₁ [° C.] in annealing is 0.1 to 20 ° C./s. The lower limit of the temperature at which the rate of temperature increase is 20 ° C./s or less is set to (Ac ₁ [° C.] − 100 ° C.) or more. The lower limit of the recrystallization temperature of the DP steel of the present invention is (Ac ₁ [° C.]). This is because it becomes −100 ° C. or higher. In addition, although the recrystallization temperature of the steel of the present invention is increased by the inclusion of one or both of Ti and Nb, it does not exceed the Ac ₁ transformation point. Further, the upper limit of the temperature of the heating rate and 20 ° C. / s or less 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 exceeds 20 ° C./s, the progress of recrystallization is remarkably suppressed. Therefore, the area ratio of non-recrystallized ferrite is remarkably increased, resulting in deterioration of ductility. Furthermore, since recrystallization temperature becomes high in steel with a high content of Nb and Ti, recrystallization hardly proceeds. When producing a steel having a small difference between the recrystallization temperature and the Ac ₁ transformation temperature, it is preferable to set the rate of temperature rise to 10 ° C./s or less in order to secure unrecrystallized ferrite.

  On the other hand, if the rate of temperature rise is slower than 0.1 ° C./s, recrystallization promotes and unrecrystallized ferrite cannot be secured, so the lower limit is made 0.1 ° C./s or more. Further, in the case of continuous annealing, it is necessary to slow the sheet feeding speed in order to slow the temperature rising rate. Therefore, from the viewpoint of productivity, the temperature rising rate is preferably 1 ° C./s or more.

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)
... (Formula 2)

  Next, the reasons for limiting the steel components in the present invention will be described.

  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 if it is less than 0.01%, the manufacturing cost increases, so the lower limit is 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 Si content exceeds 1.50%, weldability deteriorates and non-plating properties can be induced, so the upper limit is made 1.50%.

  Mn has an effect of increasing strength as an element contributing to solid solution strengthening, and is also effective for obtaining a hard second phase, and an appropriate amount is added according to a target strength level. If the Mn content is less than 0.5%, it is difficult to obtain high strength, so the lower limit is made 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%.

  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%.

  Nb and Ti are the most important elements in the present invention, and one or both of them are added. By adding Nb and Ti, fine carbides such as NbC and TiC are precipitated, and the yield strength is increased. Further, the fine precipitates can suppress recrystallization of the processed ferrite in the annealing process after cold rolling, and can promote the residual of non-recrystallized ferrite. In order to obtain such an effect, it is preferable to add one or both of Nb and Ti at a lower limit of 0.005% or more. On the other hand, one or both of Nb and Ti, when the content exceeds 0.100%, recrystallization is remarkably suppressed, and ductility may be lowered. Therefore, the upper limit of each is 0.100% or less. It is preferable to do. When the total content of one or both of Nb and Ti is 0.130% or more, recrystallization is remarkably suppressed and ductility may be lowered. Therefore, the upper limit may be made less than 0.130%. preferable.

  Furthermore, you may contain 1 type, or 2 or more types of Mo, B, Cr, Ni.

  Mo, B, Cr, and Ni are all elements that enhance hardenability, and one or more of them may be added as necessary. In order to obtain the effect of improving the strength, each of Mo: 0.1% or more, B: 0.0005% or more, Cr: 0.10% or more, Ni: 0.10% or more may be added as lower limits. preferable. On the other hand, excessive addition leads to an increase in alloy costs, so the upper limit of each is Mo: 1.5% or less, B: 0.0100% or less, Cr: 1.50% or less, Ni: 1.50% or less It is preferable that

Incidentally, when the Ac ₁ is less than 700 ° C., Nb, the recrystallization temperature progression becomes recrystallization high temperatures is significantly suppressed than Ac ₁ by the addition of Ti. Thereby, the area ratio of non-recrystallized ferrite becomes excessive, and ductility is lowered. Therefore, it is necessary that Ac ₁ is 700 ° C. or higher in the present invention.

  The microstructure of the steel sheet obtained by the present invention is composed of ferrite and a hard second phase, and 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 is a hard phase and 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%. In addition, since the non-recrystallized ferrite maintains the texture after cold rolling, increasing the area ratio of the non-recrystallized ferrite having {112} <110> orientation is effective in improving the Young's modulus. It is.

  Non-recrystallized ferrite and other ferrites, that is, recrystallized ferrite and transformed ferrite, are crystal orientation measurement data (Electron Back Scattering Pattern, EBSP) measured 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 Young's modulus is improved by the development of the {112} <110> orientation in the ½ layer thickness of the steel of the present invention. Therefore, it is preferable that the pole density of {112} <110> in the ½ layer thickness is 6 or more. Since the Young's modulus increases with an increase in the pole density of {112} <110>, it is preferably 8 or more, more preferably 10 or more.

  The pole density is synonymous with the X-ray random intensity ratio, and is the ratio of the X-ray intensity of the test material based on the X-ray intensity of a standard sample that does not accumulate in a specific orientation. The extreme density is obtained by measuring the X-ray intensity of the standard sample and the test material by the X-ray diffraction method under the same conditions and dividing the X-ray intensity of the obtained test material by the X-ray intensity of the standard sample. It is done. Further, instead of the X-ray diffraction method, a statistically sufficient number of measurements may be performed by an EBSP method or an ECP (Electron Channeling Pattern) method.

  The pole density in the {112} <110> orientation is determined by the series expansion method using a plurality of pole figures among the {110}, {100}, {211}, {310} pole figures measured by X-ray diffraction. What is necessary is just to obtain | require from the calculated three-dimensional texture (ODF). That is, the pole density in the {112} <110> orientation is represented by the X-ray random intensity ratio of (112) [1-10] in the φ2 = 45 ° section of the ODF.

  Samples to be subjected to X-ray diffraction are thinned to a predetermined plate thickness by mechanical polishing or the like so that the half-thickness surface becomes the measurement surface, and then strain is removed by chemical polishing or electrolytic polishing. Make it. When measurement inconvenience occurs, for example, when a segregation zone or a defect exists in the center thickness layer of the steel plate, an appropriate surface becomes the measurement surface in the range of 3/8 to 5/8 of the plate thickness. It is sufficient to adjust the sample.

  Here, {hkl} <uvw> indicates that the normal direction of the plate surface of the X-ray sample collected by the above method is parallel to {hkl} and the rolling direction is parallel to <uvw>. ing. 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, since the present invention is intended for the body-centered cubic structure, for example, (111), (-111), (1-11), (11-1), (-1-11), (-11-1) ), (1-1-1), and (-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). In the present invention, however, [hkl] (uvw) ) And {hkl} <uvw> are synonymous.

  Next, a manufacturing method and its preferable conditions will be described.

  The steel piece to be subjected to hot rolling may be manufactured by a conventional method, and the steel may be melted and cast. From the viewpoint of productivity, continuous casting is preferable, and it may be manufactured with a thin slab caster or the like. Further, a process such as continuous casting-direct rolling in which hot rolling is performed immediately after casting may be used. Hot rolling may be performed by a conventional method, and conditions such as rolling temperature, rolling reduction, cooling rate, and winding temperature are not particularly specified. After hot rolling, the steel sheet is cold-rolled and annealed to obtain a cold-rolled steel sheet.

However, in the present invention, in order to improve the Young's modulus and increase the rigidity in addition to workability and impact resistance, the temperature of the finish rolling and the range from 950 ° C. to the finish rolling temperature in the hot rolling process. The total rolling reduction is important. When the temperature of the finish rolling is less than the Ar ₃ transformation point, rolling is performed in the α + γ two-phase region, so that a texture unfavorable for improving the Young's modulus during cold rolling may develop and the Young's modulus may decrease.

Further, the total rolling reduction within the range from 950 ° C. to the finish rolling temperature is preferably increased in order to obtain a cold-rolled steel sheet having developed {112} <110> orientation effective for improvement of Young's modulus. This is because when the austenite phase of Ar ₃ transformation temperature to 950 ° C. is stable and rolled in a low temperature range, recrystallization due to the introduced strain hardly occurs and a hot rolled texture develops. When such a hot-rolled texture is developed, crystal rotation occurs by subsequent cold rolling, and the {112} <110> orientation develops. Furthermore, if recrystallization is suppressed during the annealing of the cold-rolled sheet, a cold-rolled steel sheet having a high Young's modulus can be obtained.

  In order to obtain the effect of improving the Young's modulus, the total rolling reduction within the range from 950 ° C. to the finish rolling temperature is preferably 30% or more. Thereby, the {112} <110> orientation develops sufficiently during cold rolling. The total rolling reduction within the range from 950 ° C. to the finish rolling temperature is a percentage obtained by dividing the difference between the plate thickness at 950 ° C. and the plate thickness after finish rolling by the plate thickness at 950 ° C. It is a representation.

Ar ₃ [° C.] is the content of C, Si, P, Al, Mn, Mo, Cu, Cr, and Ni, expressed in mass%, respectively (% C), (% Si), (% P), ( What is necessary is just to calculate by the following formula | equation using (% Al), (% Mn), (% Mo), (% Cu), (% Cr), (% Ni). In addition, the selectively added elements Mo, Cu, Cr, and Ni may be calculated as 0 when the content is about an impurity.
Ar ₃ [° C.] = 901-325 (% C) +33 (% Si) +287 (% P) +40 (% Al) −92 {(% Mn) + (% Mo) + (% Cu)} − 46 {( % Cr) + (% Ni)} (Formula 3)

  When producing the steel sheet of the present invention having excellent workability and impact resistance, the rolling reduction ratio of cold rolling is not particularly specified. However, at a cold rolling ratio of less than 10%, it is difficult to control the sheet thickness and the shape is poor. For this reason, the lower limit is preferably 10% or more. On the other hand, if the cold rolling rate exceeds 90%, the load on the rolling roll increases, and it is necessary to increase the heating rate of annealing in order to promote recrystallization and secure unrecrystallized ferrite. become. Therefore, the upper limit of the cold rolling reduction is preferably 90% or less.

  However, in the present invention, in order to improve Young's modulus and increase rigidity in addition to workability and impact resistance, it is preferable to perform cold rolling at a rolling reduction of more than 60%. This is because the {112} <110> orientation effective for improving the Young's modulus can be developed by performing cold rolling at a high rolling reduction exceeding 60%.

In the present invention, annealing after cold rolling is extremely important, and it is necessary to carry out under the above-mentioned conditions. 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 applied 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. This is because the alloying does not proceed sufficiently below 450 ° C, and the alloying proceeds excessively above 600 ° C, the plating layer becomes brittle, and the plating peels off by processing such as pressing. This is because it may trigger. 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 pieces obtained by melting and casting steel having the composition shown in Table 1 were reheated at 1250 ° C. and then hot-rolled according to a conventional method. At this time, the finishing temperature was 900 ° C., and the winding temperature was 600 ° C. Then, after performing cold rolling at a reduction rate of 60%, annealing was performed under the conditions shown in Table 2. Table 1 also shows the calculated values of Ac ₁ [° C.] and Ac ₃ [° C.]. In addition, [-] in Table 1 means that no component is intentionally added. Nb + Ti is the total amount of Nb and Ti, and [−] was calculated as 0. 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.].

  Among the cold-rolled steel sheets shown in Table 2, production No. About 3 and 6, after an annealing process, after being immersed in Zn plating bath, manufacture No.3. 6 was further alloyed at 500 ° C. for 20 s. Furthermore, among the cold-rolled steel sheets shown in Table 2, the production No. As for No. 9, it was cooled from the soaking temperature to 300 ° C. at the cooling rate of 50 ° C./s as described above, and after performing an overaging treatment at 400 ° C. for 400 s, it was cooled to room temperature at 10 ° C./s.

  From the cold-rolled steel sheet after production, a No. 5 tensile test piece of JIS Z 2201 was taken with the width direction (referred to as the TD direction) as the longitudinal direction, and the tensile characteristics in the TD direction were evaluated according to JIS Z 2241. t-El [%] is the elongation at break, and L-El [%] is the local elongation, which is a value obtained by subtracting the elongation at break from the elongation at maximum force, that is, the uniform elongation.

  Microstructure observation of the plate thickness section of the steel sheet was performed with an optical microscope using a sample taken with the rolling direction as the observation surface and etching as a repeller method. The area ratio of the hard second phase was obtained as a total of phases other than ferrite by image analysis of a structure photograph taken with an optical microscope. Further, the area ratio and the balance of the non-recrystallized ferrite, that is, the area ratio of the ferrite excluding the non-recrystallized ferrite were obtained by image analysis by collating the crystal orientation measurement of EBSP and the measurement result with the optical micrograph.

  The results are shown in Table 3. In the present invention, the product of tensile strength TS [MPa] which is an index of ductility and total elongation t-EL [%], tensile strength TS [MPa] which is an index of stretch flangeability and local elongation L-El [%] ], The ratio of the yield strength YS [MPa] to the tensile strength TS [MPa], which is an index of impact resistance characteristics, that is, TS × t-El [MPa ·%], TS × L-El [MPa ·%] and Those having YP / TS × 100 [%] of 18000 [MPa ·%], 7000 [MPa ·%] and 80 [%] or more were evaluated as good.

  As shown in Table 3, the steel having the chemical composition of the present invention is hot-rolled and cold-rolled under appropriate conditions, and further annealed under appropriate conditions. It is possible to obtain a high-strength cold-rolled steel sheet that is excellent in strength-ductility balance, stretch flangeability, and impact resistance properties even when subjected to a heat treatment.

On the other hand, Steel No. Since J has a small amount of C, steel no. Since K has a small amount of Mn, the strength is lowered, and TS × t-El [MPa ·%] is also lowered. Steel No. L has a low Ac ₁ transformation temperature. Since M has a high Nb and Ti content, recrystallization hardly progresses, and unrecrystallized ferrite remains excessively. Therefore, both total elongation and local elongation are reduced, and TS × t-El [MPa ·%] and TS × L-El [MPa ·%] are also reduced. Steel No. N has a low content of Nb and Ti, so that the yield strength is reduced, and recrystallization proceeds to reduce unrecrystallized ferrite, resulting in a decrease in local elongation, YR [%] and TS × L-El [ [MPa ·%] is also decreased.

In addition, production No. In _No. 2, the rate of temperature increase from (Ac ₁ [° C.]-100 ° C.) to Ac ₁ [° C.] is fast, excessive unrecrystallized ferrite remains, the strength is low, and the total elongation and local elongation both decrease. TS × t-El [MPa ·%] and TS × L-El [MPa ·%] are also decreased.

Production No. No. 5 has a high maximum temperature for annealing, so _No. 8 has a long residence time at Ac ₁ [° C.] or higher, so that there is little unrecrystallized ferrite, the hard second phase is increased, and the strength is high, but the total elongation and local elongation are reduced, and TS × t-El [MPa ·%] and TS × L-El [MPa ·%] are also decreased.

  Production No. No. 11 has a low maximum ultimate temperature of annealing and a hard second phase was not obtained, so the strength is low, and TS × t-El [MPa ·%] and TS × L-El [MPa ·%] are also reduced. Yes.

Steel having the composition shown in Table 4 was melted in a vacuum melting furnace, hot-rolled under the conditions shown in Table 5, wound up and pickled at 600 ° C., and then cold-rolled and annealed under the conditions shown in Table 5. Went. The average cooling rate from Ac ₁ [° C.] to 500 ° C. or the overaging temperature was 50 ° C./s. Here, [-] in Table 4 means that no component is intentionally added. Table 4 also shows the calculated values of Ac ₁ [° C.] and Ac ₃ [° C.].

In Table 5, the rolling reduction in the hot rolling step is 950 ° C. or less and the total rolling reduction until finish rolling, and was determined from the plate thickness at 950 ° C. and the plate thickness after finish rolling. Further, FT [° C.] is a hot rolling finishing temperature. 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 No. 19 and 23, after an annealing process, after being immersed in Zn plating bath, manufacture No. For No. 23, alloying treatment was further performed at 500 ° C. for 20 s. Furthermore, among the cold-rolled steel sheets shown in Table 5, the production No. As for No. 29, it was cooled from a soaking temperature to 300 ° C. at a cooling rate of 50 ° C./s as described above, and after performing an overaging treatment for 400 s at 300 ° C., it was cooled to room temperature at 10 ° C./s.

  The tensile properties in the TD direction of the cold-rolled steel sheet after production were evaluated in the same manner as in Example 1. Further, the microstructure observation of the plate thickness cross section of the steel plate and the crystal orientation measurement of EBSP were performed in the same manner as in Example 1, and the area ratio of the hard second phase, the area ratio of the non-recrystallized ferrite, and the non-recrystallized ferrite were measured. The area ratio of the excluded ferrite was determined. The pole density in the {112} <110> orientation in the 1/2 plate thickness part was measured by the X-ray diffraction method. A sample for X-ray diffraction was prepared by mechanical polishing and electrolytic polishing so that the half-thick surface was the measurement surface.

The Young's modulus was measured by performing the transverse resonance method described in JIS Z 2280 at room temperature. That is, the Young's modulus was calculated from the following equation by applying vibration without fixing the sample, measuring the primary resonance frequency by gradually changing the frequency of the oscillator.
E = 0.946 × (l / h) ³ × m / w × f ²
Here, E: dynamic Young's modulus [N / m ² ], l: length of the test piece [m], h: thickness of the test piece [m], m: mass of the test piece [kg], w: Width of test piece [m], f: primary resonance frequency [s ⁻¹ ] of lateral resonance method.

  The results are shown in Table 6. As in Example 1, TS × t-El [MPa ·%], TS × L-El [MPa ·%], and YP / TS × 100 [%] are 18000 [MPa ·%] and 7000 [MPa, respectively. -%] And 80 [%] or more were evaluated as good. Further, the Young's modulus E [GPa], which is an index of rigidity, was evaluated as good when the modulus was 240 [GPa] or more.

  As shown in Table 6, by appropriately annealing the steel having the chemical composition of the present invention, it is a high-strength cold steel that has excellent workability and impact resistance characteristics even if it is over-aged, Zn plated, or alloyed. It is possible to obtain a rolled steel sheet. Furthermore, the steel having the chemical composition of the present invention is hot-rolled and cold-rolled under appropriate conditions, and further annealed under appropriate conditions, so that it has high strength with excellent rigidity in addition to workability and impact resistance. It is possible to obtain a cold-rolled steel sheet. Production No. No. 22 has a low rolling reduction in the hot rolling process, so No. 25 has a high finishing temperature, so No. 28 is an example in which the E [GPa] is lowered because the texture is not sufficiently developed because the cold rolling rate is low.

On the other hand, Steel No. Since AJ has a small amount of C, steel no. Since AK has less Mn, the strength is reduced and TS × t-El [MPa ·%] is also reduced. Steel No. AL has a low Ac ₁ transformation temperature. Since AM has a high Nb and Ti content, recrystallization hardly proceeds and unrecrystallized ferrite remains excessively. Therefore, both total elongation and local elongation are reduced, and TS × t-El [MPa ·%] and TS × L-El [MPa ·%] are also reduced. Steel No. AN has a low Nb and Ti content, and therefore yield strength is reduced, recrystallization proceeds and unrecrystallized ferrite is reduced, local elongation is reduced, and YR [%] and TS × L-El [ [MPa ·%] is also decreased.

  In addition, production No. No. 31 is an example in which the rate of temperature increase in the annealing process is high and the unrecrystallized ferrite remains excessively, the strength is low, both the total elongation and the local elongation decrease, and TS × t-El [MPa ·%] and TS × L-El [MPa ·%] is also decreased. Production No. No. 34 has a high maximum temperature for annealing, so that there is little unrecrystallized ferrite, the hard second phase is increased, and although it is high strength, the total elongation and local elongation are reduced, and TS × t-El [MPa %] And TS × L-El [MPa ·%] are also decreased. Production No. No. 38 has a short holding time at the highest temperature of annealing, and a sufficient hard second phase was not obtained. Therefore, the strength is low, and TS × t-El [MPa ·%] and TS × L-El [MPa · %] Is also decreasing.

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%,
Si: 0.01 to 1.50%,
Mn: 0.50 to 3.50%
P: 0.150% or less,
S: 0.0150% or less,
Al: 0.200% or less,
N: 0.0100% or less, and
Nb: 0.005 to 0.100%,
Ti: 0.005 to 0.100%
One or both of the above is contained in a total of less than 0.130%, the balance is made of iron and inevitable impurities, Ac ₁ [° C.] is 700 ° C. or higher, and the metal structure is made of ferrite and a hard second phase, The ferrite is composed of one or both of recrystallized ferrite and transformed ferrite and non-recrystallized ferrite, the area ratio of the non-recrystallized ferrite is 10 to 70%, and the area of one or both of the recrystallized ferrite and transformed ferrite A high-strength cold-rolled steel sheet excellent in local ductility, workability, and impact resistance, wherein the rate is 10 to 70% and the area ratio of the hard second phase is 1 to 30%.
Here, Ac ₁ [° C.] is the Ac ₁ transformation obtained from the following formula (Formula 1) by the contents (% C), (% Mn), and (% Si) of C, Mn, and Si expressed in mass%. Temperature.
_{Ac 1 = 761.3 + 212 (%} C) -45.8 (% Mn) +16.7 (% Si)
... (Formula 1) Furthermore, in 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 having excellent local ductility, workability, and impact resistance characteristics according to claim 1, comprising one or more of them. 3. The local ductility, workability, and impact resistance characteristics according to claim 1, wherein the pole density in the {112} <110> orientation in the ½ layer thickness is 6 or more. High strength cold-rolled steel sheet with excellent resistance. A high-strength cold-rolled steel sheet excellent in local ductility, workability and impact resistance characteristics, characterized in that hot-dip Zn plating is provided on the surface of the cold-rolled steel sheet according to any one of claims 1 to 3. A high-strength cold-rolled steel sheet excellent in local ductility, workability, and impact resistance characteristics, characterized in that alloyed hot-dip Zn plating is provided on the surface of the cold-rolled steel sheet according to any one of claims 1 to 3. . A steel slab having the chemical composition according to claim 1 or 2 is hot-rolled, pickled, and cold-rolled, and then the steel plate is (Ac ₁ [° C] -100 ° C). Ac ₁ [℃] ~ a heating rate of up to Ac ₁ [° C.] as 0.1 to 20 ° C. / s from _{{Ac 1 [℃] + 2} /3 × (Ac 3 [℃] -Ac 1 [℃]) }, The steel sheet is annealed with a residence time within the temperature range of 10 to 200 s, and is excellent in local ductility, workability and impact resistance. A method for producing a cold-rolled steel sheet.
Here, Ac ₁ [° C.] and Ac ₃ [° C.] are expressed by mass% of C, Mn, and Si (% C), (% Mn), and (% Si) according to the following (formula 1) and These are the Ac ₁ transformation temperature and Ac ₃ transformation temperature determined from the 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)
... (Formula 2) The steel slab having the chemical composition according to any one of claims 1 and 2, wherein the finishing rolling temperature is not less than the Ar ₃ transformation temperature, and the total rolling reduction within the range from 950 ° C to the finishing rolling temperature is 30%. 7. Hot ducting as described above, pickling, cold rolling at a reduction rate of more than 60%, and annealing the steel sheet , local ductility, workability and impact resistance characteristics For producing high-strength cold-rolled steel sheets with excellent resistance. After annealing according to any one of claims 6 and 7, the steel is cooled to 350 to 500 ° C, and then hot-dip Zn plating is applied, and high strength excellent in local ductility, workability and impact resistance characteristics A method for producing a cold-rolled steel sheet. A high-strength cold-rolled steel sheet excellent in local ductility, workability, and impact resistance characteristics, characterized by performing heat treatment for 10 seconds or more in a temperature range of 450 to 600 ° C after performing hot-dip Zn plating according to claim 8. Production method. The local ductility, workability, and impact resistance characteristics are characterized by subjecting the cold-rolled steel sheet produced by the method according to any one of claims 6 to 9 to 0.1% to 5.0% skin pass rolling. An excellent method for producing high-strength cold-rolled steel sheets.

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