JP2005120471A - High strength steel sheet manufacturing method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high strength steel sheet manufacturing method which is useful for applications to steel sheets for cars or the like, has high tensile strength (TS) of ≥ 440 MPa with a high r-value (r-value ≥ 1.2) and with an absolute value of Δr being small as 0.3 or under. <P>SOLUTION: In this manufacturing method, a steel slab having a composition consisting of, by mass, 0.010-0.050% C, ≤ 1.0% Si, 1.0-3.0% Mn, 0.005-0.1% P, ≤ 0.01% S, 0.005-0.5% Al, ≤ 0.01% N, and 0.01-0.12% Nb while Nb content and C content satisfy a relationship of (Nb/93)/(C/12) = 0.2-0.7 is subjected to the finish hot roll at the finish rolling outlet side temperature of ≥ 800°C. The rolled steel slab is coiled at a coiling temperature of 400-720°C to form a hot-rolled sheet. The hot-rolled steel sheet is cold-rolled after the pickling to form a cold-rolled sheet. The cold-rolled sheet is annealed within a predetermined temperature range, and then, cooled at an average cooling rate of ≥ 5°C/s in the temperature range from the annealing temperature to 500°C. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention is a high-strength steel sheet having a tensile strength (TS) of 440 MPa or more, and has a high Rankford value (average r value ≧ 1.2), which is useful for use in steel sheets for automobiles, and has an r value. It is an object of the present invention to propose an advantageous method for producing a high-strength steel sheet having a small absolute value of in-plane anisotropy (Δr).

In recent years, in order to regulate CO ₂ emissions from the viewpoint of global environmental conservation, there has been a demand for improved fuel efficiency of automobiles. In addition, in order to ensure the safety of passengers in the event of a collision, it is also required to improve safety centered on the collision characteristics of the automobile body. Thus, both weight reduction and reinforcement of the automobile body are being actively promoted.

  In order to satisfy the weight reduction and strengthening of automobile bodies at the same time, it is said that it is effective to increase the strength of component materials within a range where there is no problem with rigidity and to reduce the weight by reducing the plate thickness. (Also known as high tension.) Steel plates are actively used in automotive parts.

  Since the weight reduction effect increases as the strength of the steel sheet used increases, the automotive industry tends to use steel sheets having a tensile strength (TS) of 440 MPa or more as panel materials for inner and outer plates, for example.

  On the other hand, since many automobile parts made of steel plates are formed by press working, the steel plates for automobiles are required to have excellent press formability. However, the high-strength steel sheet is greatly deteriorated in formability, particularly deep drawability, as compared with a normal mild steel sheet. Therefore, TS ≧ 440 MPa, more preferably TS ≧ 500 MPa, More preferably, TS ≧ 590 MPa, and there is an increasing demand for steel sheets having good deep drawability, and the average r value is a Rankford value (hereinafter referred to as “r value”) which is an evaluation index of deep drawability. A high-strength steel sheet of ≧ 1.2 is required. Further, it has been found that when the absolute value of the in-plane anisotropy of the r value becomes small, for example 0.3 or less, the formability in the actual press becomes better even if the average r-value level is the same. A small anisotropy is also required.

As a means to increase the strength while having a high r value, the base is based on IF (interstitial atom free) steel added with Ti and Nb in amounts that fix carbon and nitrogen that dissolve in the ultra-low carbon steel sheet. As such, there is a method of adding a solid solution strengthening element such as Si, Mn, P or the like, for example, a method disclosed in Patent Document 1.
JP-A-56-139654

Patent Document 1 has a composition of C: 0.002 to 0.015%, Nb: C% x 3 to C% x 8 + 0.020%, Si: 1.2% or less, Mn: 0.04 to 0.8%, P: 0.03 to 0.10% This is a technique related to a high-tensile cold-rolled steel sheet having excellent formability and non-aging properties of a tensile strength of 35 to 45 kg / mm class ² (340 to 440 MPa class).

  However, the technology of adding solid solution strengthening elements using such ultra-low carbon steel as a raw material produces high-strength steel sheets with a tensile strength (also called tensile strength) of 440 MPa or higher, or even 500 MPa or higher and 590 MPa or higher. Attempts have been made to increase the amount of alloy elements added, resulting in problems such as surface appearance problems, deterioration of plating properties, and manifestation of secondary work brittleness. Further, when a solid solution strengthening component is added in a large amount, the r value deteriorates, so that there is a problem that the level of the r value decreases as the strength is increased.

Further, in order to reduce the amount of C to an extremely low carbon region of less than 0.010% as specifically disclosed in the above cited reference 1, vacuum degassing must be performed in the steelmaking process, that is, A large amount of CO ₂ is generated in the manufacturing process, and it is difficult to say that this technique is preferable from the viewpoint of global environmental conservation.

  As a method for increasing the strength of a steel sheet, there is a structure strengthening method other than the solid solution strengthening method as described above. For example, there is a DP (Dual-Phase) steel sheet, which is a composite structure steel sheet composed of a soft ferrite phase and a hard martensite phase. DP steel sheet is generally good in terms of ductility, has an excellent strength-ductility balance (TS × El), and has a low yield ratio, that is, yield stress is low for tensile strength, Although it is characterized by excellent shape freezing properties during press molding, it has a low r value and inferior deep drawability. This is said to be because solute C, which is essential for martensite formation, inhibits the formation of {111} recrystallized texture effective for increasing the r value.

  As an attempt to improve the r value of such a composite structure steel plate, for example, there is a technique of Patent Document 2 or Patent Document 3.

Patent Document 2 discloses a method in which after cold rolling, box annealing is performed at a temperature from the recrystallization temperature to the Ac ₃ transformation point, and then heating to 700 to 800 ° C. to obtain a composite structure, followed by quenching and tempering. It is disclosed. However, in this method, since the quenching and tempering is performed at the time of continuous annealing, the manufacturing cost becomes a problem. Further, box annealing is inferior in terms of processing time and efficiency as compared to continuous annealing.
Japanese Patent Publication No.55-10650 JP 55-100934

  In the technique of Patent Document 3, after cold rolling to obtain a high r value, box annealing is first performed, the temperature at this time is set to a two-phase region of ferrite (α) -austenite (γ), and then continuous annealing is performed. Is what you do. In this technique, Mn is concentrated from the α phase to the γ phase during soaking of the box annealing. This Mn-concentrated phase preferentially becomes a γ phase during the subsequent continuous annealing, and a mixed structure can be obtained even at a cooling rate of the order of a gas jet. However, this method requires a relatively high temperature and long-time box annealing to concentrate Mn, and therefore, in the manufacturing process, such as frequent adhesion between steel plates, generation of temper collar, and decrease in the life of the furnace inner cover. There are many problems.

Patent Document 4 discloses that C: 0.003-0.03%, Si: 0.2-1%, Mn: 0.3-1.5%, Ti: 0.02-0.2% (however, (effective Ti) / (C + N) atomic concentration ratio) 0.4 to 0.8), which is hot-rolled, cold-rolled, and then subjected to continuous annealing that is rapidly cooled after heating to a predetermined temperature, and is a composite structure type excellent in deep drawability and shape freezeability A method for producing a high-tensile cold-rolled steel sheet is disclosed. Specifically, a steel having a composition of 0.012% C-0.32% Si-0.53% Mn-0.03% P-0.051% Ti in cold mass is cold-rolled. After heating to 870 ° C, which is a two-phase region of α-γ, and then cooling at an average cooling rate of 100 ° C / s, a composite structure type cold rolled steel sheet with r value = 1.61 and TS = 482 MPa can be manufactured. The effect is shown. However, in order to obtain a high cooling rate of 100 ° C / s, water quenching equipment is required, and water-quenched steel sheets have surface treatment problems, so there are problems with manufacturing equipment and materials. is there.
Japanese Patent Publication No. 1-35900

Furthermore, Patent Document 5 discloses a technique for improving the r value of a composite structure steel sheet by optimizing the V content in relation to the C content. This is because by precipitating C in the steel as V-type carbides before recrystallization annealing to reduce the amount of dissolved C as much as possible to achieve a high r value, and subsequently heating in the two-phase region of α-γ, V-type carbides are dissolved to concentrate C in γ, and a martensite phase is generated in the subsequent cooling process. However, the addition of V causes an increase in cost because it is expensive, and VC precipitated in the hot-rolled sheet increases the deformation resistance during cold rolling, so that it is disclosed in the examples. Cold rolling at a rolling reduction of 70% has a manufacturing problem such as increasing the load on the rolls to increase the risk of trouble occurrence and concern about a decrease in productivity. There is also a concern that the recrystallization behavior changes due to the influence of fine precipitates and the in-plane anisotropy of the r value increases.
Japanese Patent Laid-Open No. 2002-226941

Moreover, there exists a technique of patent document 6 as a technique of the high strength steel plate excellent in deep drawability, and its manufacturing method. This technique obtains a high-strength steel sheet containing a predetermined amount of C, having an average r value of 1.3 or more, and having a total of 3% or more of one or more of bainite, martensite, and austenite in the structure. The manufacturing method includes a cold rolling reduction rate of 30 to 95%, followed by annealing to increase the r value by developing a texture by forming Al and N clusters and precipitates, and subsequently in the structure. In addition, heat treatment is performed so as to have at least 3% of at least one of bainite, martensite, and austenite. In this method, after cold rolling, annealing for obtaining a good r value and heat treatment for forming a structure are required, and in the annealing process, the holding time is a long time of 1 hour or more. There is a problem that productivity is poor in terms of process (time). Furthermore, since the second phase fraction of the obtained structure is relatively high, it is difficult to stably secure an excellent strength ductility balance.
JP 2003-64444 A

  In order to increase the strength of a (soft) steel sheet excellent in deep drawability, a method for increasing the strength by solid solution strengthening that has been conventionally studied requires the addition of a large amount or an excess of alloy components. In terms of cost, process, and improvement of the r value, there are problems.

  In addition, the method using the strengthening of the structure requires a two-time annealing (heating) method and a high-speed cooling facility, and thus has a problem in the manufacturing process. Further, although a method using VC is disclosed, it is expensive. Addition of V causes an increase in cost, and precipitation of VC increases deformation resistance during rolling, which also makes stable production difficult.

  The present invention proposes a method for producing a high-strength steel sheet having a good average draw value and a good deep drawability with a small average r value and a small in-plane anisotropy of the r value. For the purpose.

  The present invention has been intensively studied to solve the above-mentioned problems, and the C content range is set to 0.010 to 0.050 mass%, an Nb amount corresponding to the C content is added, and cold rolling is further performed. After performing, it discovered that the in-plane anisotropy of a high r value and r value could be made small by performing annealing in the optimal range according to the amount of components.

  That is, the lower limit of the annealing temperature needs to be T1 = 800 + 625 Nb according to the amount of Nb that changes the recrystallization temperature in order to cause recrystallization.

  In addition, the upper limit of the annealing temperature requires a transformation temperature constraint, and the two-phase structure during annealing must be controlled. For this reason, it discovered that it was necessary to optimize the upper limit of annealing temperature to the temperature matched with the amount of Mn etc. which change a transformation point. In this case, since the change in the γ fraction during heating is complicated and it is difficult to accurately measure the transformation point, the average r value and the r The in-plane anisotropy of the value was evaluated, and the upper limit of the annealing temperature at which the average r value ≧ 1.2 and the absolute value of the in-plane anisotropy of the r value ≦ 0.3 was obtained as a function of the component contained.

  And it determined that the in-plane anisotropy of the high r value and the r value can be reduced by setting the lower limit and the upper limit of the annealing temperature in this way and performing the annealing within the optimum range according to the amount of the contained component.

That is, the gist of the present invention is as follows.
(I) In mass%,
C: 0.010 to 0.050%
Si: 1.0% or less
Mn: 1.0-3.0%
P: 0.005-0.1%
S: 0.01% or less
Al: 0.005-0.5%
N: 0.01% or less and
Nb: 0.01-0.12%
And Nb content and C content are
(Nb / 93) / (C / 12) = 0.2-0.7
(Nb and C in the formula are the contents of each element (mass%))
A steel slab having a composition satisfying the above relationship is subjected to finish rolling at a finish rolling temperature of 800 ° C or higher by hot rolling, and wound at a winding temperature of 400 to 720 ° C to form a hot rolled sheet A rolling process, a cold rolling process in which the hot-rolled sheet is pickled and cold-rolled to obtain a cold-rolled sheet, and an annealing temperature for the cold-rolled sheet is expressed by the following formulas (1) and (2) T1 A high-strength steel sheet characterized by having a cold-rolled sheet annealing step of annealing at an average cooling rate in the temperature range from the annealing temperature to 500 ° C: 5 ° C / s or more. Production method.
T1 （℃） ＝ 800 ＋ 625Nb ・ ・ ・ ・ (1)
T2 (° C) = 950-45 (Mn-Si-5P) (2)
Here, the element symbol in a formula shows content (mass%) of each element.

(II)% by mass
C: 0.010 to 0.050%
Si: 1.0% or less
Mn: 1.0-3.0%
P: 0.005-0.1%
S: 0.01% or less
Al: 0.005-0.5%
N: 0.01% or less and
Nb: 0.01-0.12%
And Nb content and C content are
(Nb / 93) / (C / 12) = 0.2-0.7
(Nb and C in the formula are the contents of each element (mass%))
A hot-rolling process in which a steel slab having a composition satisfying the relationship is hot-rolled to form a hot-rolled sheet having an average crystal grain size of 8 μm or less, and the hot-rolled sheet is cold-rolled, In the cold rolling process, the annealing temperature of the cold-rolled sheet is set to T1 ° C. or higher and T2 ° C. or lower as shown in the following formulas (1) and (2), and then in the temperature range from the annealing temperature to 500 ° C. A method for producing a high-strength steel sheet, comprising: an average cooling rate: a cold-rolled sheet annealing step for cooling at 5 ° C./s or more.
T1 （℃） ＝ 800 ＋ 625Nb ・ ・ ・ ・ (1)
T2 (° C) = 950-45 (Mn-Si-5P) (2)
Here, the element symbol in a formula shows content (mass%) of each element.

(III) In addition to the above composition, the steel slab is further selected from Mo: 0.02 mass% or more, Cr: 0.05 mass% or more, Cu: 0.05 mass% or more, and Ni: 0.05 mass% or more, or The above (I) is characterized in that it contains 0.5% by mass or less in total of two or more types, and the upper limit of the annealing temperature is annealed at T2 ° C. shown in the following formula (3) instead of the formula (2). Or the manufacturing method of the high strength steel plate as described in (II).
T2 (° C) = 950-45 (Mn-Si-5P-0.8Mo + Cu + Ni) (3)
Here, the element symbol in a formula shows content (mass%) of each element.

(IV) In addition to the above composition, the steel slab further contains Ti: 0.1% by mass or less, and the contents of Ti, S and N in the steel are
(Ti / 48) / {(S / 32) + (N / 14)} ≦ 2.0 (Ti, S and N in the formula are the contents of each element (mass%))
The method for producing a high-strength steel sheet according to the above (I), (II) or (III), wherein the following relationship is satisfied.

(V) The high-strength steel plate according to any one of (I) to (IV) above, further comprising a plating treatment step of forming a plating layer on the surface of the steel plate after the cold-rolled plate annealing step. Manufacturing method.

  This invention is necessary for the formation of martensite when the C content is in the range of 0.010 to 0.050% by mass without thorough reduction of solid solution C that adversely affects deep drawability like conventional ultra-low carbon IF steels. Despite the state in which a certain amount of solute C remains, a {111} recrystallized texture is developed to secure an average r value ≧ 1.2 and good deep drawability, and the steel structure is made of ferrite. By forming a composite structure having a phase and a second phase including a martensite phase, high strength of TS440 MPa or more, more preferably TS500 MPa or more is achieved. Furthermore, in this invention, since the in-plane anisotropy of the r value can be reduced, particularly excellent press formability can be ensured.

Although the reason for this is not necessarily clear, it can be considered as follows.
In order to increase the r value, that is, to develop the {111} recrystallization texture, in the conventional mild steel sheet, the solute C before cold rolling and recrystallization is reduced as much as possible and the hot rolled sheet structure is refined. To do so has been an effective means. On the other hand, in the DP steel sheet as described above, since the solid solution C necessary for martensite formation is required, the recrystallization texture of the parent phase does not develop and the r value is low. However, in the present invention, it has been newly found that there is an exquisite good component range that enables both the {111} recrystallized texture development of the ferrite phase as a parent phase and the formation of the martensite phase. That is, the C content is 0.010 to 0.050% by mass, in which the C content is lower than that of the conventional DP steel sheet (low carbon steel level) and the C content is higher than that of the ultra-low carbon steel. It was newly found that by adding Nb appropriately according to the amount, both {111} recrystallization texture development and martensite phase formation can be achieved simultaneously.

  As conventionally known, since Nb has a recrystallization delay effect, it is possible to refine the hot-rolled sheet structure by appropriately controlling the finishing temperature during hot rolling. Nb has a high carbide forming ability.

  In the present invention, in particular, the hot rolling finish temperature is set to an appropriate range immediately above the transformation point, and in addition to refining the hot rolled sheet structure, the coil winding temperature after hot rolling is also set appropriately so that hot rolling NbC is precipitated in the plate to reduce solid solution C before cold rolling and before recrystallization.

  Here, by setting the Nb content and the C content to satisfy (Nb / 93) / (C / 12) = 0.2 to 0.7, C that does not precipitate as NbC is present.

  Conventionally, it has been said that the presence of such C inhibits the development of {111} recrystallized texture. However, in the present invention, the total C content is not precipitated and fixed as NbC, but is necessary for the formation of the martensite phase. A high r-value can be achieved while molten C is present.

  The reason for this is not clear, but rather than a negative factor for the formation of {111} recrystallized texture due to the presence of solute C, in addition to refinement of the hot rolled sheet structure, fine NbC is precipitated in the matrix. This is probably because the positive factor of accumulating strain near the grain boundary during cold rolling and promoting the generation of {111} recrystallized grains from the grain boundary is greater. In particular, the effect of precipitating NbC in the matrix is not effective at the C content of the conventional ultra-low carbon steel, and is not effective for the first time in the proper range (0.010 to 0.050% by mass) of the C content of the present invention. It is presumed to be exhibited, and finding the appropriate range of the C content is the basis of the technical idea of the present invention.

  C other than NbC and its existence form are presumed to be cementite-based carbides or solute C, but the presence of C that is not fixed as NbC can form a martensite phase during cooling in the annealing process. And succeeded in increasing the strength.

  According to the manufacturing method of the present invention, compared to the prior art, a degassing step for making an ultra-low carbon steel is not necessary in the steel making step, and addition of an excessive alloy element for utilizing solid solution strengthening. Is also unnecessary and advantageous in terms of cost. Furthermore, there is no need to add a special element such as V which increases the alloy cost and rolling load.

The present invention is described in detail below.
The unit of element content is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.
First, the reason for limiting the component composition of the steel slab used in the present invention will be described.

C: 0.010 to 0.050%
C is an important element in the present invention together with Nb described later. C is effective for increasing the strength and promotes the formation of a composite structure having a ferrite phase as a main phase and a second phase including a martensite phase. Therefore, in the present invention, C is 0.010% or more from the viewpoint of forming a composite structure. It is necessary to contain, preferably 0.015% or more. On the other hand, if the C content exceeds 0.050%, a good r value cannot be obtained, so the upper limit of the C content is 0.050%, preferably 0.030% or less.

Si: 1.0% or less
Si promotes ferrite transformation and raises the C content in untransformed austenite, making it easier to form a composite structure of ferrite and martensite phases, and has the effect of solid solution strengthening. If the content exceeds V, surface defects called red scales are generated during hot rolling, which deteriorates the surface appearance of the steel sheet. In addition, when hot dip galvanizing is performed, the wettability of the plating is deteriorated and plating unevenness occurs, and the plating quality deteriorates. Therefore, the Si content must be 1.0% or less, preferably 0.7%. The following. In order to obtain the above effect, the lower limit of the Si content is preferably 0.01%, more preferably 0.05%.

Mn: 1.0-3.0%
Mn is effective in increasing strength and has the effect of lowering the critical cooling rate at which a martensite phase is obtained, and promotes the formation of the martensite phase during cooling after annealing, so the required strength level and after annealing The Mn is preferably an element effective for preventing hot cracking due to S. From such a viewpoint, Mn needs to be contained at 1.0% or more, preferably 1.2% or more. On the other hand, containing excessive Mn exceeding 3.0% degrades the r value and weldability, so the upper limit of the Mn content is 3.0%.

P: 0.005-0.1%
P is an element having an effect of solid solution strengthening. However, if the P content is less than 0.005%, not only the effect does not appear, but also the dephosphorization cost increases in the steel making process. Therefore, P is contained in an amount of 0.005% or more, preferably 0.01% or more. On the other hand, if the P content exceeds 0.1%, P segregates at the grain boundaries, and the secondary work brittleness resistance and weldability deteriorate. Moreover, when it is set as the hot dip galvanized steel sheet, the diffusion of Fe from the steel sheet to the plated layer at the interface between the plated layer and the steel sheet is suppressed during the alloying process after the hot dip galvanizing, and the alloying processability is deteriorated. For this reason, an alloying treatment at a high temperature is required, and the obtained plating layer is likely to undergo plating peeling such as powdering and chipping. Therefore, the upper limit of the P content is set to 0.1%.

S: 0.01% or less S is an impurity and causes hot cracking, and also exists as inclusions in steel and deteriorates various properties of the steel sheet. Therefore, it is necessary to reduce it as much as possible. Specifically, the S content is 0.01% or less because it is acceptable up to 0.01%.

Al: 0.005-0.5%
In addition to being useful as a deoxidizing element for steel, Al has the effect of fixing solid solution N contained as an impurity to improve normal temperature aging resistance. Must be 0.005% or more. On the other hand, the Al content exceeding 0.5% causes high alloy costs and further induces surface defects, so the upper limit of Al content is set to 0.5%. More preferably, it is 0.1% or less.

N: 0.01% or less N is an element that degrades aging resistance at room temperature, and is an element that is preferably reduced as much as possible. If the N content exceeds 0.01%, a large amount of Al or Ti needs to be added to prevent deterioration of normal temperature aging resistance due to N, so the upper limit of the N content is set to 0.01%.

Nb: 0.01 to 0.12% and (Nb / 93) / (C / 12) = 0.2 to 0.7
Nb is the most important element in the present invention, and has the effect of refining the hot-rolled sheet structure and precipitating and fixing C as NbC in the hot-rolled sheet, and contributes to increasing the r value. From such a viewpoint, Nb needs to be contained in an amount of 0.01% or more, preferably 0.03% or more. On the other hand, excessive Nb content exceeding 0.12% makes it difficult to control the in-plane anisotropy and tends to increase the in-plane anisotropy of the r value, so the upper limit of the Nb content is set to 0.12%. .

  In order to achieve the effect of Nb content, the Nb content (mass%) and the C content (mass%) should satisfy the relational expression of (Nb / 93) / (C / 12) = 0.2 to 0.7. It is necessary to make it a range. Here, the element symbol (Nb, C) in the formula indicates the content (% by mass) of each element. (Nb / 93) / (C / 12) represents the atomic concentration ratio of Nb and C. When (Nb / 93) / (C / 12) is less than 0.2, the amount of solid solution C is large, and the formation of {111} recrystallized texture effective for increasing the r value is inhibited. Further, if (Nb / 93) / (C / 12) exceeds 0.7, the amount of C necessary to form the martensite phase is prevented from being present in the steel, so the martensite phase is finally included. A structure having a second phase cannot be obtained. Therefore, the Nb content is set to 0.01 to 0.12%, and the Nb content and the C content are set to satisfy the range of (Nb / 93) / (C / 12) = 0.2 to 0.7. More preferably, the range satisfies (Nb / 93) / (C / 12) = 0.2 to 0.5.

The above is the basic composition of the high-strength steel sheet of the present invention.
In the present invention, in addition to the above-described composition, one or more selected from Mo, Cr, Cu and Ni shown below and / or Ti may be added as necessary. Good.

Mo: 0.02% or more, Cr: 0.05% or more, Cu: 0.05% or more and Ni: 0.05% or more selected from a total of 0.5% by mass or less
Mo, Cr, Cu and Ni, like Mn, have the effect of slowing the critical cooling rate at which a martensite phase is obtained, and are elements that promote the formation of the martensite phase during cooling after annealing, and are effective in improving the strength level. There is. Mo, Cr, Cu and Ni may be contained in the steel as unavoidable impurities in the range of Mo: less than 0.02%, Cr: less than 0.05%, Cu: less than 0.05% and Ni: less than 0.05%. In order to obtain the effect, the Mo content is 0.02% or more, Cr: 0.05% or more, Cu: 0.05% or more, and Ni: 0.05% or more. It is preferable to make it. However, excessive addition exceeding 0.5% in total of one or more elements of these components not only saturates the effect but also increases the cost due to expensive components, so Mo, Cr The upper limit of the total content of one or more elements of Cu and Ni is preferably 0.5%. The lower limit values for the contents of Mo, Cr, Cu and Ni are more preferably 0.05%, 0.1%, 0.1% and 0.1%, respectively, and the upper limit values for the respective contents of Mo, Cr, Cu and Ni. Is preferably 0.5%.

Ti: 0.1% or less, and the contents of Ti, S, and N in the steel are (Ti / 48) / {(S / 32) + (N / 14)} ≦ 2.0
Ti is an element which is equivalent to Al or more effective than Al and is effective in precipitation fixation of solute N. In order to obtain this effect, Ti is preferably contained in an amount of 0.005% or more. However, excessive addition exceeding 0.1% not only increases the cost, but also prevents the solid solution C necessary for the formation of the martensite phase from remaining in the steel due to the formation of TiC. Therefore, the Ti content is preferably 0.1% or less.

  Further, Ti preferentially bonds with S and N in the steel, and then bonds with C. Considering the decrease in Ti yield due to the formation of inclusions in steel, etc., when Ti addition amount exceeds (Ti / 48) / {(S / 32) + (N / 14)} 2.0, S and N The effect of Ti addition of fixing is saturated, and on the contrary, the effect of preventing the formation of TiC and preventing the solid solution C from remaining in the steel becomes large. Therefore, the Ti content should satisfy (Ti / 48) / {(S / 32) + (N / 14)} ≦ 2.0 in relation to the contents of S and N that are preferentially bonded in steel. Is preferred. In addition, Ti, S, and N in this relational expression here are content (mass%) of each element.

  In the present invention, the balance other than the above-described components is preferably substantially composed of iron and inevitable impurities. Inevitable impurities include, for example, Mo: less than 0.02%, Cr: less than 0.05%, Cu: less than 0.05% and Ni: less than 0.05%, and Ti: less than 0.005%, as described above.

  In addition, if it is in the normal steel composition range, even if it contains B, Ca, REM, etc., there is no problem. For example, B is an element having an effect of improving the hardenability of steel and can be contained as necessary. However, when the content exceeds 0.003%, the effect is saturated, and therefore it is preferably 0.003% or less.

  Moreover, Ca and REM have the effect | action which controls the form of a sulfide type inclusion, and, thereby, prevents the deterioration of the various characteristics of a steel plate. Since such an effect tends to be saturated when the content of one or two selected from Ca and REM exceeds 0.01% in total, it is preferable to make the content less than this.

  Other inevitable impurities include, for example, Sb, Sn, Zn, Co, etc. The allowable ranges of these contents are Sb: 0.01% or less, Sn: 0.1% or less, Zn: 0.01% or less , Co: 0.1% or less.

The high-strength steel sheet produced by the method of the present invention satisfies the above component composition and steel structure, and satisfies an average r value of 1.2 or more.
Here, the “average r value” means an average plastic strain ratio obtained by JIS Z 2254, and is a value calculated from the following formula.
Average r value = (r ₀ + 2r ₄₅ + r ₉₀ ) / 4
R ₀ , r _45, and r ₉₀ are plastic strain ratios obtained by measuring test pieces in the 0 °, 45 °, and 90 ° directions, respectively, with respect to the rolling direction of the plate surface.

In addition, the high-strength steel sheet produced by the method of the present invention satisfies the above-described components, steel microstructure and properties, and is parallel to the plate surface as a texture determined by X-ray diffraction at a steel plate 1/4 thickness position. The integral intensity ratios P ₍₂₂₂₎ , P ₍₂₀₀₎ , P ₍₁₁₀₎ , and P ₍₃₁₀₎ of the (222) plane, (200) plane, (110) plane, and (310) plane are P _{(222) / {P (200) + P} (110) + P (310)} is preferable to satisfy ≧ 1.5, more preferably from _{P (222) / {P (} 200) + P (110) + P (310)} ≧ 2.0 .
FIG. 1 shows the calculation of the r value and the value of P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ } for various manufactured steel sheets of the present invention and comparative steel sheets, and based on these calculated values. Is plotted.

Conventionally, the r value is high when the plate surface has a texture parallel to the {111} plane, but the r value is known to be low in the texture parallel to the {110} plane or the {100} plane.
As a result of diligent research on the correlation between the r value and the texture in the steel sheet of the present invention, the details are not clear yet, but the {310} face is small, but the r value is lowered as in the {100} and {110} faces. It was a texture, and it was found that reducing this contributes to a higher r-value. This is not clear in detail, but the Nb addition has a high rolling reduction in the unrecrystallized γ region during hot rolling, the fine NbC precipitation described above, and the presence of C that is not precipitated and fixed as NbC. It is thought that it contributes to {310} plane reduction.

In addition, {111} texture means that the <111> direction of the crystal is oriented in the direction perpendicular to the steel plate surface. From the crystallographic and Bragg reflection conditions, in the case of α-Fe with a body-centered cubic structure, diffraction on the {111} plane does not occur on the (111) plane but occurs on the (222) plane, so X-ray diffraction The value of (222) plane (P ₍₂₂₂₎ ) was used as the integral intensity ratio. The (222) plane is substantially the same as the <111> direction because the [222] direction is oriented in the direction perpendicular to the steel plate face. Therefore, a high intensity ratio of the (222) plane corresponds to the development of {111} texture. For the same reason, the value of (200) plane (P ₍₂₀₀₎ ) was used for {100} plane.

  Here, the X-ray diffraction integrated intensity ratio is a relative intensity based on the X-ray diffraction integrated intensity of a non-directional standard sample (irregular sample). The X-ray diffraction may be either an angle dispersion type or an energy dispersion type, and the X-ray source may be a characteristic X-ray or a white X-ray. It is desirable to measure 7 to 10 measurement surfaces (110) to (420) which are the main diffraction surfaces of α-Fe. Moreover, the steel plate 1/4 plate thickness position specifically refers to a range of 1/8 to 3/8 of the plate thickness of the steel plate measured from the steel plate surface, and X-ray diffraction is an arbitrary value within this range. You can do it in terms of

  In the present invention, a hot rolling process in which a steel slab having a composition in the above-described range is hot-rolled to form a hot-rolled sheet, and a cold-rolled sheet by cold-rolling the hot-rolled sheet by pickling and cold rolling. A high-strength steel sheet can be produced by sequentially performing a cold rolling process and a cold-rolled sheet annealing process for annealing the cold-rolled sheet.

  In the present invention, the steel slab is first subjected to finish rolling at a finish rolling exit temperature of 800 ° C. or higher by hot rolling, and wound at a winding temperature of 400 to 720 ° C. to obtain a hot rolled sheet (hot Rolling process).

  The steel slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macro segregation of components, but may be produced by an ingot-making method or a thin slab casting method. Moreover, after manufacturing the steel slab, in addition to the conventional method of once cooling to room temperature and then heating again, the steel slab is charged directly into the heating furnace without being cooled and charged in a heating furnace. Energy-saving processes such as direct feed rolling and direct rolling, in which hot rolling is performed immediately after heat insulation, can also be applied without problems.

  Here, the slab heating temperature is preferably low because the precipitates are coarsened to develop a {111} recrystallized texture and improve deep drawability. However, if the heating temperature is less than 1000 ° C., the rolling load increases and the risk of trouble during hot rolling increases, so the slab heating temperature is preferably 1000 ° C. or higher. Note that the upper limit of the slab heating temperature is preferably set to 1300 ° C. because of an increase in scale loss accompanying an increase in oxidized weight.

  The steel slab heated under the above conditions is subjected to hot rolling for rough rolling and finish rolling. Here, the steel slab is made into a sheet bar by rough rolling. The conditions for rough rolling need not be specified, and may be determined according to a conventional method. From the viewpoint of lowering the slab heating temperature and preventing troubles during hot rolling, it is preferable to use a so-called sheet bar heater that heats the sheet bar.

  Next, the sheet bar is finish-rolled to obtain a hot-rolled sheet. At this time, the finish rolling outlet temperature (FT) is 800 ° C. or higher. This is to obtain a hot-rolled sheet structure capable of obtaining excellent deep drawability after cold rolling and annealing. If the FT is less than 800 ° C., the structure has a processed structure, and not only the {111} texture does not develop after cold rolling annealing, but also the rolling load during hot rolling becomes high. Therefore, the finish rolling outlet temperature (FT) is 800 ° C. or higher. When FT exceeds 980 ° C., the structure becomes coarse, which also prevents formation and development of {111} recrystallized texture after cold rolling annealing, which is disadvantageous in obtaining a high r value. The upper limit is preferably 980 ° C.

  Moreover, in order to reduce the rolling load at the time of hot rolling, it is good also as lubrication rolling between some or all passes of finishing rolling. Lubrication rolling is effective from the viewpoint of uniform steel plate shape and uniform material. The coefficient of friction during lubrication rolling is preferably in the range of 0.10 to 0.25. Furthermore, it is also preferable to use a continuous rolling process in which the adjacent sheet bars are joined and finish-rolled continuously. The application of the continuous rolling process is also desirable from the viewpoint of the operational stability of hot rolling.

  The coil winding temperature (CT) is in the range of 400 to 720 ° C. This temperature range is an appropriate temperature range for depositing NbC in the hot-rolled sheet. When CT exceeds 720 ° C., the crystal grains become coarse and hinders the increase in r value after cold rolling annealing. On the other hand, when the CT is lower than 400 ° C., the precipitation of NbC is difficult to occur, which is disadvantageous for increasing the r value. The coil winding temperature (CT) is preferably 550 to 680 ° C.

By performing the hot rolling step, a hot rolled steel sheet having an average crystal grain size of 8 μm or less can be obtained. That is, the high-strength steel sheet of the present invention is a hot-rolled sheet having a composition within the above-described range and having an average crystal grain size of 8 μm or less, and cold-rolled by subjecting the hot-rolled sheet to cold rolling. It can manufacture by passing through the cold-rolling process to perform, and the cold-rolled sheet annealing process which achieves recrystallization and composite organization to this cold-rolled sheet sequentially.
In conventional mild steel sheets, it is known that the r value increases as the crystal grain size of the hot-rolled sheet is refined.

2 (a), 2 (b), 3 (a), and 3 (b) are optical micrographs of a hot rolled steel sheet that has undergone nital corrosion. Nital solution is 3% nitric acid alcohol solution _{(3% HNO 3 -C 2 H} 5 0H), was 10~15s corrosion.

  Here, FIG. 2 (a) is 0.033% C and Nb-free, the average crystal grain size of the hot rolled sheet: 8.9 μm, the average r value of the steel sheet obtained by cold rolling annealing: 0.9, FIG. (B) is 0.035% C-0.015% Nb {(Nb / 93) / (C / 12)} = 0.06}, obtained by subjecting the hot rolled sheet to an average crystal grain size of 5.9 μm and cold rolling annealing. The average r value of the steel sheet: 1.0, FIG. 3 (a) is 0.035% C-0.083% Nb {(Nb / 93) / (C / 12)} = 0.31}, and the average grain size of the hot rolled sheet : 5.6 μm, average r value of steel sheet obtained by cold rolling annealing: 1.3, FIG. 3B is 0.035% C−0.072% Nb {(Nb / 93) / (C / 12)} = 0.27} The average crystal grain size of the hot-rolled sheet is 2.8 μm, the average r value of the steel sheet obtained by cold-rolling annealing is 1.5, and FIGS. 3A and 3B show the component composition within the range of the present invention. It is a hot-rolled steel sheet.

  FIG. 2 (a) shows Nb-free steel that deviates from the steel of the present invention in terms of the components. The hot rolled sheet has an average crystal grain size of 8 μm or more and a low r value. In FIG. 2 (b), although the hot-rolled sheet structure is refined by adding Nb, the value of (Nb / 93) / (C / 12) is out of the scope of the present invention, so the effect is exhibited. The r value is low. In FIGS. 3A and 3B, the value of (Nb / 93) / (C / 12) is within the range of the present invention, the hot-rolled sheet structure is refined, and the r value is increased.

In the hot-rolled sheet structure, Nb addition causes a line (1) that is deeply corroded by the nital liquid as a grain boundary as well as a line (2) that is shallowly corroded.
In the present invention, when measuring the grain size, the crystal grain size was measured using the above-mentioned lines (1) and (2) as grain boundaries.

  In general, the crystal grain size is often referred to as a so-called large-angle grain boundary having an inclination of 15 ° or more and a so-called small-angle grain boundary having a tilt angle of less than 15 °. As a result of EBSP (Electron Back Scatter Diffraction Pattern) analysis of the shallow-corrosion line (2), it was found that the shallow-corrosion line (2) is a so-called small-angle grain boundary with an inclination of less than 15 °. In the present invention, the hot-rolled sheet is characterized by a large number of so-called small-inclined grain boundaries, that is, the above-described line (2), which has an inclination of less than 15 °. As a result of measuring the grain size with both of (2) as the grain boundary, when the average crystal grain size exceeds 8 μm, the effect of increasing the r value of the high strength steel sheet of the present invention does not appear, and the average grain size It has been found that by reducing the diameter to 8 μm or less, an effect can be obtained in increasing the r value to an average r value of 1.2 or more. Therefore, the average crystal grain size of the hot-rolled sheet is 8 μm or less.

As a result of EBSP analysis of the structure of the steel of the present invention, the measurement of the crystal grain size using the above-mentioned lines (1) and (2) as grain boundaries means that the grain boundaries having an inclination of 5 ° or more are grain boundaries. It was confirmed that this was equivalent to measuring the particle size as a boundary.
Further, from this, although details are not clear, it is presumed that an inclination angle of 5 ° or more is effective in promoting the generation of recrystallization nuclei preferable for deep drawability from the grain boundary in the present invention.

As a method for measuring the crystal grain size, a microscopic structure is imaged using an optical microscope for a plate thickness cross section (L cross section) parallel to the rolling direction, and a cutting method according to JIS G 0552 or ASTM is performed on the sample surface. The average intercept length l (μm) of the crystal grain at (ASTM) nominal grain diameter d _n = 1.13 × l may be found, and the average grain diameter may be obtained using a device such as EBSP. Also good.

  In the present invention, the section length of the average particle diameter was obtained by cutting a microstructure according to JISG0552 on a cross section of the plate thickness parallel to the rolling direction by imaging the microstructure with an optical microscope. That is, using the microscopic microstructure photographed, the number of ferrite crystal grains cut at a certain length of the line segment in the rolling direction and the direction perpendicular to the rolling direction according to JISG0552 is measured, and the length of the line segment is measured. Is divided by the number of ferrite crystal grains to be cut by the line segment, and is obtained as an intercept length in each direction, and an average (arithmetic mean) value of these values is obtained as an average intercept length l of the crystal grains here. (Μm).

  Further, as described above, by adjusting the component composition and hot rolling conditions, 15% or more of the total C content can be precipitated and fixed as NbC in the hot-rolled sheet stage, that is, precipitated and fixed as NbC. The ratio of the C content to the total C content in the steel can be 15% or more, which is advantageous for increasing the r value.

The ratio of the amount of C that is precipitated and fixed as NbC to the total amount of C in the steel (hereinafter simply referred to as “the ratio of the amount of C that is fixed and fixed”) refers to the chemical analysis (extraction analysis) of the hot-rolled sheet. It is a value calculated by the following formula from the amount of precipitated Nb obtained.
[C] _fix = 100 x 12 x ([Nb ^* ] / 93) / [C] _total
Here, when Ti is not contained in the steel, Nb forms NbN,
[Nb ^* ] = [Nb] − (93 [N] / 14), [Nb ^* ]> 0
On the other hand, when Ti is contained in the steel, N preferentially forms TiN, so [Nb ^* ] = [Nb] − (93 [N ^* ] / 14)
In the formula,
[N ^* ] = [N] − (14 [Ti ^* ] / 48), [N ^* ]> 0
[Ti ^* ] = [Ti] − (48 [S] / 32), [Ti ^* ]> 0
[C] _fix is the ratio (%) of the amount of C fixed by precipitation,
[C] _total is the total C content (mass%) in the steel,
[Nb], [N], [Ti], and [S] are precipitation Nb, precipitation N, precipitation Ti, and precipitation S amount (mass%), respectively.

  Reducing solute C at the stage before cold rolling and recrystallization is effective for increasing the r value, and the amount of C precipitated and fixed as NbC is 15% or more of the total C content. The effect appears preferably at 20% or more, and this can be achieved by the production method of the present invention.

  Next, the hot-rolled sheet is cold-rolled to obtain a cold-rolled sheet (cold-rolling step). Here, the hot-rolled sheet is preferably pickled before cold rolling in order to remove scale. What is necessary is just to carry out on normal conditions as pickling conditions. The cold rolling condition is not particularly limited as long as it can be a cold-rolled sheet having a desired size and shape, but the rolling reduction during cold rolling is preferably at least 40%, more preferably 50% or more. And A high cold rolling reduction ratio is generally effective for increasing the r value. If the reduction ratio is less than 40%, a {111} recrystallized texture does not develop, and it becomes difficult to obtain excellent deep drawability. On the other hand, in this invention, the r value increases as the cold rolling reduction is increased in the range up to 90%, but when it exceeds 90%, not only the effect is saturated, but also the load on the roll during cold rolling is increased. Therefore, the upper limit is preferably 90%.

Next, the cold rolled sheet is annealed at an annealing temperature of T1 ° C. or higher and T2 ° C. or lower as shown below, and then cooled at an average cooling rate in the temperature range from the annealing temperature to 500 ° C .: 5 ° C./s or higher (cool Sheet annealing process).
T1 （℃） ＝ 800 ＋ 625Nb
T2 (° C) = 950-45 (Mn-Si-5P)
Here, the element symbol in a formula shows content (mass%) of each element.
However, one or more of Mo, Cr, Cu, and Ni are added to the steel slab, that is, Mo: 0.02 mass% or more, Cr: 0.05 mass% or more, Cu: 0.05 mass% or more, and Ni: 0.05 mass% or more In the case of containing one or more selected from the above, the upper limit of the annealing temperature is set to T2 ° C. shown in the following formula (3) instead of the formula (2).
T2 (° C) = 950-45 (Mn-Si-5P-0.8Mo + Cu + Ni) (3)
Here, the element symbol in a formula shows content (mass%) of each element.

  The cold-rolled sheet annealing step is preferably performed in a continuous annealing line (CAL) or a continuous hot-dip galvanizing line (CGL) in order to ensure the cooling rate required in the present invention.

  When the annealing temperature is lower than T1 (° C.), the carbide does not dissolve so that the second martensite phase can be sufficiently formed, and as a result, a sufficient balance between strength and elongation cannot be obtained. Further, since Nb is added, if the annealing temperature is less than T1 (° C.), the crystal grains are elongated in the rolling direction even after recrystallization, and the r-value in-plane anisotropy and elongation are deteriorated. Therefore, by setting the annealing temperature to T1 (° C.) or higher, the carbide can be sufficiently dissolved, and the crystal grains can have a shape preferable for the in-plane anisotropy and elongation of the r value.

  On the other hand, when the annealing temperature is higher than T2 (° C), the crystal grains become coarse, the {111} recrystallized texture does not easily develop, and the deep drawability tends to deteriorate, and the cause is unknown. The in-plane anisotropy of the r value becomes worse. Furthermore, the martensite phase of the second phase becomes excessive, and the strength-elongation balance characteristic is remarkably deteriorated. In addition, as described above, T2 (° C.) is not only an equilibrium transformation point, but various experiments were conducted by heat treatment simulating continuous annealing with steels having different components. The upper limit of the annealing temperature at which the average r value ≧ 1.2 and the absolute value of the in-plane anisotropy of the r value can be 0.3 or less is expressed as a function of the component based on the result of investigation of the internal anisotropy. .

  From the viewpoint of forming a martensite phase, the cooling rate after the annealing needs to be cooled by setting the average cooling rate in the temperature range from the annealing temperature to 500 ° C. to 5 ° C./s or more. If the average cooling rate in the temperature range is less than 5 ° C./s, the martensite phase is difficult to form and a ferrite single phase structure is formed, resulting in insufficient structure strengthening.

The steel sheet produced according to the present invention preferably forms a second phase including a martensite phase, and the average cooling rate up to 500 ° C. is required to be equal to or higher than the critical cooling rate. In order to achieve this, This is satisfied by setting the average cooling rate in the temperature range from the annealing temperature to 500 ° C. to 5 ° C./s or more. Although there is no particular limitation on cooling at less than 500 ° C, it is preferable to continue cooling to 300 ° C at an average cooling rate of 5 ° C / s or higher, and when overaging is performed, up to the overaging temperature. The average cooling rate is preferably 5 ° C./s or more.
Note that the upper limit of the cooling rate is not particularly required from the viewpoint of martensite phase formation, and cooling may be performed using water quenching equipment or the like in addition to roll cooling or gas jet cooling.

  Moreover, after the said cold-rolled sheet annealing process, surface treatments, such as an electroplating process or a hot dipping process, may be given and a plating layer may be formed in the steel plate surface.

For example, when performing hot dip galvanizing treatment, which is often used for automotive steel sheets, as the plating treatment, the cold rolled sheet annealing step is performed in a continuous hot dip plating line, followed by cooling after the annealing in a hot dip galvanizing bath. It suffices to form a hot dip galvanized layer on the surface. In this case, after the hot dip galvanized solution is melted, it is preferably cooled so that the average cooling rate up to 300 ° C. is 5 ° C./s or more. . Moreover, after immersing in the hot dip galvanizing bath following the cooling after the annealing to form a hot dip galvanized layer on the surface, an alloying treatment may be further performed to manufacture an alloyed hot dip galvanized steel sheet. In this case, in the cooling after the alloying treatment, it is preferable that the average cooling rate up to 300 ° C. is 5 ° C./s or more.
In addition, regarding the cooling after leaving the hot dip galvanizing bath or after the alloying treatment, the upper limit of the cooling rate is not particularly required from the viewpoint of martensite phase formation. Alternatively, cooling may be performed using water quenching equipment or the like.

  Further, the cooling after the annealing may be performed in the annealing line, and after cooling to room temperature, the hot dip galvanizing may be separately performed in the hot dip galvanizing line, or further alloying treatment may be performed.

  Here, the plating layer is not limited to pure zinc and zinc-based alloy plating, but may of course be various plating layers conventionally applied to the steel sheet surface, such as pure Al and Al-based alloy plating.

  Further, the steel sheet produced as described above may be subjected to temper rolling or leveler processing for the purpose of adjusting the shape correction, surface roughness, and the like. The total elongation of temper rolling or leveler processing is preferably in the range of 0.2 to 15%. If it is less than 0.2%, the intended purpose of shape correction and roughness adjustment may not be achieved. On the other hand, if it exceeds 15%, it tends to cause a significant decrease in ductility. In addition, although the processing form differs between temper rolling and leveler processing, it has been confirmed that there is no significant difference in the effect between the two. Temper rolling and leveler processing are effective even after plating.

  The high-strength steel plate manufactured by the manufacturing method of the present invention described above has a steel structure including a ferrite phase with an area ratio of 50% or more and a martensite phase with an area ratio of 1% or more, and an average r value of 1.2 or more. In addition, the absolute value of the in-plane anisotropy of the r value is as small as 0.3 or less.

Here, the “average r value” means an average plastic strain ratio obtained by JIS Z 2254, and is a value calculated from the following formula.
Average r value = (r ₀ + 2r ₄₅ + r ₉₀ ) / 4

R ₀ , r _45, and r ₉₀ are plastic strain ratios obtained by measuring test pieces in the 0 °, 45 °, and 90 ° directions, respectively, with respect to the rolling direction of the plate surface.

Further, the in-plane anisotropy of the r value means the in-plane anisotropy similarly obtained by JIS Z 2254, and is a value calculated from the following formula.
In-plane anisotropy Δr = (r ₀ −2r ₄₅ + r ₉₀ ) / 2

  Below, the steel structure of the high strength steel plate obtained by the manufacturing method of this invention is demonstrated.

(1) Having a steel structure containing a ferrite phase with an area ratio of 50% or more and a martensite phase with an area ratio of 1% or more The high-strength steel sheet obtained by the production method of the present invention has an area ratio of 50% or more. It is preferable to use a composite structure steel plate having a steel structure including a ferrite phase of 1% and a martensite phase with an area ratio of 1% or more. Here, a {111} recrystallization texture of the ferrite phase occupying an area ratio of 50% or more is developed, and an average r value ≧ 1.2 is achieved.

  Such a high-strength steel sheet has a good deep drawability and includes a ferrite phase having an area ratio of 50% or more and a martensite phase having an area ratio of 1% or more in order to obtain a steel sheet having a tensile strength ≧ 440 MPa. Has a steel structure. When the ferrite phase is reduced and the area ratio is less than 50%, it becomes difficult to ensure good deep drawability, and the press formability tends to decrease. The ferrite phase is preferably 70% or more in terms of area ratio, and in order to utilize the advantages of the composite structure, the ferrite phase is preferably 99% or less in area ratio.

  Here, the “ferrite phase” includes a bainitic ferrite phase having a high dislocation density transformed from an austenite phase in addition to a polygonal ferrite phase.

  In the present invention, it is necessary that a martensite phase is present, and the martensite phase is contained in an area ratio of 1% or more. When the martensite phase is less than 1%, it is difficult to obtain a good strength ductility balance. Furthermore, the martensite phase is preferably 3% or more. The martensite phase is preferably 30% or less in order to improve the γ value.

  In addition, a structure including a pearlite phase, a bainite phase, or a retained austenite (γ) phase in addition to the above-described ferrite phase and martensite phase may be used. Furthermore, in order to sufficiently obtain the effects of the ferrite phase and the martensite phase, it is preferable that the total of the area ratio of the ferrite phase and the area ratio of the martensite phase is 80% or more. The ferrite phase and the martensite phase may be identified by observation with a scanning electron microscope, and the area ratio may be obtained by imaging a microscopic structure with a scanning electron microscope.

(2) The average r value is 1.2 or more and the absolute value of Δr is 0.3 or less. The high-strength steel sheet of the present invention satisfies the same component composition and steel structure as the steel slab and the average r value ≧ 1.2. In addition, the absolute value of the in-plane anisotropy Δr of the r value satisfies 0.3 or less. In the present invention, the composition is adjusted to the above component composition, and the steel structure is manufactured so as to have a steel structure containing a ferrite phase and a martensite phase. For the first time, the average r value is 1.2 or more and the absolute value of Δr is 0.3 or less. Could be achieved.

  The high-strength steel plate obtained by the production method of the present invention includes a cold-rolled steel plate, a steel plate having a plating layer after being subjected to surface treatment such as electroplating or hot dip galvanizing, so-called plated steel plate. Here, “plating” refers to pure zinc, zinc-based alloy plating in which zinc is the main component and alloy elements are added, or in addition to pure Al, Al-based alloy plating in which Al is the main component and alloy elements are added. In addition, a plating layer conventionally applied to the surface of a steel sheet is also included.

Next, examples of the present invention will be described.
Molten steel having the composition shown in Table 1 was melted in a converter and made into a slab by a continuous casting method. These steel slabs were heated to 1250 ° C. and roughly rolled into sheet bars, and then hot-rolled sheets were formed by a hot rolling process in which finish rolling under the conditions shown in Table 2 was performed. These hot-rolled sheets were made into cold-rolled sheets by a cold rolling process in which cold rolling with a rolling reduction of 65% was performed after pickling. Subsequently, these cold-rolled sheets were subjected to continuous annealing under the conditions shown in Table 2 in a continuous annealing line (CAL) to obtain cold-rolled steel sheets (also referred to as cold-rolled annealed sheets). Moreover, about some cold rolled sheets, it replaced with the continuous annealing line, and performed continuous annealing on the conditions shown in Table 2 in the continuous hot-dip galvanizing line (CGL). In addition, about the steel plate which performed the cold rolled sheet annealing process in the continuous hot dip galvanizing line, hot dip galvanization (plating bath temperature: 480 degreeC) was performed in-line continuously after that. Further, the obtained cold-rolled annealed sheet or hot-dip galvanized steel sheet was subjected to temper rolling with an elongation of 0.5%.

About the hot-rolled sheet after the hot rolling, the ratio of the amount of C deposited and fixed as NbC, and the results of investigation on the microstructure, tensile properties and r-value of the finally obtained cold-rolled annealed sheet Is shown in Table 2.
The survey method is as follows.

(i) Ratio of C amount deposited and fixed as NbC in hot-rolled sheet As described above, the amount of precipitated Nb, precipitated Ti, precipitated N, and precipitated S was quantified by extraction analysis, and determined by the following formula.
[C] _fix = 100 x 12 x ([Nb ^* ] / 93) / [C] _total
Here, when Ti is not contained in the steel,
[Nb ^* ] = [Nb] − (93 [N] / 14), [Nb ^* ]> 0
When containing Ti,
[Nb ^* ] = [Nb] − (93 [N ^* ] / 14)
In the formula,
[N ^* ] = [N] − (14 [Ti ^* ] / 48), [N ^* ]> 0
[Ti ^* ] = [Ti] − (48 [S] / 32), [Ti ^* ]> 0
[C] _fix is the ratio (%) of the amount of C fixed by precipitation,
[C] _total is the total C content (mass%) in the steel,
[Nb], [N], [Ti], and [S] are precipitation Nb, precipitation N, precipitation Ti, and precipitation S amount (mass%), respectively.
The extraction analysis was performed by ICP emission spectrometry after the residue obtained by electrolytic extraction using a 10% maleic acid electrolyte was alkali-melted and the melt was dissolved in acid.

(Ii) Crystal grain size of hot-rolled sheet An image of a plate thickness section (L section) parallel to the rolling direction corroded with nital is imaged with an optical microscope, and the average grain size of the average grain is as described above by a cutting method according to JIS G 0552. The section length l (μm) was determined and expressed as (ASTM) nominal particle diameter d _n = 1.13 × l. As described above, as described above, both the lines corroded by the nital solution and deeply corroded as usual and the shallowly corroded lines were counted as grain boundaries. Further, it was confirmed by EBSP analysis that the value of the average grain size measured in this way corresponds to a value measured by regarding a grain boundary having an inclination angle of 5 ° or more as a grain boundary. Here, a 3% nitric acid alcohol solution (3% HNO ₃ —C ₂ H ₅ OH) was used as the nital solution and was corroded for 10 to 15 seconds.

(iii) Microstructure of cold-rolled annealed plates Specimens were taken from each cold-rolled annealed plate, and the plate thickness cross section (L cross section) parallel to the rolling direction was fined at 1000 to 3000 times using a scanning electron microscope. The visual tissue was imaged, the types of phases were observed, the area of each phase was analyzed, and the area ratio of the ferrite phase and the area ratio of the martensite phase were determined.

(iv) Tensile properties JIS No. 5 tensile test specimens were taken from each of the obtained cold-rolled annealed plates in a direction 90 ° (C direction) with respect to the rolling direction, and the crosshead speed was 10 mm / in accordance with the provisions of JIS Z 2241. A tensile test was performed at min, and yield stress (YS), tensile strength (TS), and elongation (El) were determined.

(v) Average r value JIS No. 5 tensile test from the rolling direction (L direction) of each obtained cold rolled annealed sheet, 45 ° direction (D direction) with respect to the rolling direction, and 90 ° direction (C direction) with respect to the rolling direction. Pieces were collected. Measure the width strain and plate thickness strain of each specimen when 10% uniaxial tensile strain was applied to these specimens, and use these measurements to determine the average r value in accordance with JIS Z 2254 regulations. (Average plastic strain ratio) was calculated and used as the r value. Further, the in-plane anisotropy (Δr) of the r value was calculated in accordance with the provisions of JIS Z 2254.

(Vi) Texture The energy dispersive X-ray diffraction using white X-rays was performed at the steel plate 1/4 position of each cold-rolled annealed plate. The measurement plane is the main diffraction plane of α-Fe (110) plane, (200) plane, (211) plane, (220) plane, (310) plane, (222) plane, (321) plane, (400 ) Surface, (411) surface, (420) surface, a total of 10 surfaces, the relative intensity ratio with the non-directional standard sample to determine the X-ray diffraction integrated intensity ratio of each surface, (222) surface, Substituting the X-ray diffraction integrated intensity ratios P ₍₂₂₂₎ , P ₍₂₀₀₎ , P ₍₁₁₀₎ and P ₍₃₁₀₎ for the (200) plane, (110) plane and (310) plane into the terms on the right side of the following equation The left side term A was calculated.
A = P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ }

  As is clear from the investigation results shown in Table 2, all of the examples of the present invention have a high r value of TS440 MPa or higher, an average r value of 1.2 or higher, and the absolute value of in-plane anisotropy. Is 0.3 or less, that is, Δr = −0.3 to +0.3, which is a small value. On the other hand, in the comparative example manufactured under conditions outside the scope of the present invention, the strength is insufficient, the r value is reduced to less than 1.2, or the absolute value of Δr is greater than 0.3.

  According to the present invention, a high-strength steel sheet having a high r value of TS440 MPa or more and an average r value of 1.2 or more and a small Δr absolute value of 0.3 or less can be produced inexpensively and stably. It becomes possible and has a remarkable effect on the industry. For example, when the high-strength steel sheet of the present invention is applied to automobile parts, it is possible to increase the strength of parts that have been difficult to press-form so far, and it is possible to sufficiently contribute to collision safety and weight reduction of an automobile body. is there. Moreover, it is applicable not only to automobile parts but also to household appliance parts and pipe materials.

The average r value and the value of P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ } were calculated for various inventive steel plates produced by the method of the present invention and comparative steel plates, and these were calculated. It is the figure plotted based on the value. (A) and (b) are optical micrographs obtained by immersing a hot-rolled sheet in a nital solution to corrode the surface, and both of these values are (Nb / 93) / (C / 12). This is an example that does not satisfy the proper scope of the invention. (A) and (b) are optical micrographs obtained by immersing a hot-rolled sheet in a nital solution to corrode the surface, and both of these values are (Nb / 93) / (C / 12). It is an example satisfying the appropriate range of the invention.

Claims (5)

% By mass
C: 0.010 to 0.050%
Si: 1.0% or less
Mn: 1.0-3.0%
P: 0.005-0.1%
S: 0.01% or less
Al: 0.005-0.5%
N: 0.01% or less and
Nb: 0.01-0.12%
And Nb content and C content are
(Nb / 93) / (C / 12) = 0.2-0.7
(Nb and C in the formula are the contents of each element (mass%))
A steel slab having a composition satisfying the above relationship is subjected to finish rolling at a finish rolling temperature of 800 ° C or higher by hot rolling, and wound at a winding temperature of 400 to 720 ° C to form a hot rolled sheet A rolling process, a cold rolling process in which the hot-rolled sheet is pickled and cold-rolled to obtain a cold-rolled sheet, and an annealing temperature for the cold-rolled sheet is expressed by the following formulas (1) and (2) T1 A high-strength steel sheet characterized by having a cold-rolled sheet annealing step of annealing at an average cooling rate in the temperature range from the annealing temperature to 500 ° C: 5 ° C / s or more. Production method.
T1 （℃） ＝ 800 ＋ 625Nb ・ ・ ・ ・ (1)
T2 (° C) = 950-45 (Mn-Si-5P) (2)
Here, the element symbol in a formula shows content (mass%) of each element. % By mass
C: 0.010 to 0.050%
Si: 1.0% or less
Mn: 1.0-3.0%
P: 0.005-0.1%
S: 0.01% or less
Al: 0.005-0.5%
N: 0.01% or less and
Nb: 0.01-0.12%
And Nb content and C content are
(Nb / 93) / (C / 12) = 0.2-0.7
(Nb and C in the formula are the contents of each element (mass%))
A hot-rolling process in which a steel slab having a composition satisfying the relationship is hot-rolled to form a hot-rolled sheet having an average crystal grain size of 8 μm or less, and the hot-rolled sheet is cold-rolled, In the cold rolling process, the annealing temperature of the cold-rolled sheet is set to T1 ° C. or higher and T2 ° C. or lower as shown in the following formulas (1) and (2), and then in the temperature range from the annealing temperature to 500 ° C. A method for producing a high-strength steel sheet, comprising: an average cooling rate: a cold-rolled sheet annealing step for cooling at 5 ° C./s or more.
T1 （℃） ＝ 800 ＋ 625Nb ・ ・ ・ ・ (1)
T2 (° C) = 950-45 (Mn-Si-5P) (2)
Here, the element symbol in a formula shows content (mass%) of each element. In addition to the above composition, the steel slab is one or more selected from Mo: 0.02% by mass or more, Cr: 0.05% by mass or more, Cu: 0.05% by mass or more, and Ni: 0.05% by mass or more. 3 or less in total, and the upper limit of the annealing temperature is T2 ° C. shown in the following formula (3) instead of the above formula (2), and annealing is performed. Manufacturing method of high strength steel sheet.
T2 (° C) = 950-45 (Mn-Si-5P-0.8Mo + Cu + Ni) (3)
Here, the element symbol in a formula shows content (mass%) of each element. In addition to the above composition, the steel slab further contains Ti: 0.1% by mass or less, and the contents of Ti, S and N in the steel are
(Ti / 48) / {(S / 32) + (N / 14)} ≦ 2.0 (Ti, S and N in the formula are the contents of each element (mass%))
The method for producing a high-strength steel sheet according to claim 1, 2 or 3, wherein the following relationship is satisfied. The method for producing a high-strength steel sheet according to any one of claims 1 to 4, further comprising a plating treatment step of forming a plating layer on the surface of the steel plate after the cold-rolled sheet annealing step.