KR20060134029A - High-rigidity high-strength thin steel sheet and method for producing same - Google Patents

Abstract

Disclosed is a high-rigidity high-strength thin steel sheet having a tensile strength of not less than 590 MPa and a Young's modulus of not less than 230 GPa at the same time. The high-rigidity high-strength thin steel sheet has a composition consisting of, in mass%, C:0.02-0.15%, Si: 1.5% or less, Mn: 1.0-3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less, Ti: 0.02-0.50%, and the balance of iron and unavoidable impurities, wherein the C, N, S and Ti contents satisfy the following relations: Ti* = Ti - (47.9/14) x N - (47.9/32.1) x S >= 0.01 and 0.01 <= C - (12/47.9) x Ti* <= 0.05. The structure of such a steel sheet is mainly composed of a ferrite phase, and contains not less than 1% of a martensite phase in the area ratio.

Description

High Rigidity High Strength Steel Sheet and Manufacturing Method Thereof {HIGH-RIGIDITY HIGH-STRENGTH THIN STEEL SHEET AND METHOD FOR PRODUCING SAME}

The present invention relates to a highly rigid high strength steel sheet and a method for producing the same which are suitable for automobile bodies. The high rigidity high strength steel sheet according to the present invention is a columnar structural member whose rigidity such as a central pillar, a locker, a side frame and a cross member of an automobile has a thickness sensitivity close to 1, for applications requiring rigidity. Widely suitable.

In recent years, with increasing interest in global environmental issues, weight reduction of vehicle bodies is very important, such as regulation of exhaust gas even in automobiles. For this reason, it is effective to reduce the vehicle body weight by increasing the strength of the steel sheet while reducing the thickness of the steel sheet.

In recent years, the increase in strength of the steel sheet has proceeded remarkably, and the use of the thin steel sheet whose plate | board thickness is less than 2.0 mm increased. It is inevitable to increase the strength and to make the thickness thinner for the weight reduction of the vehicle body, while at the same time controlling the deterioration of the strength of the parts. The problem that the strength of the component is lowered by making the thickness of such a steel sheet thinner occurs in a steel sheet having a tensile strength of 590 MPa or more, and in particular, this problem becomes serious in a steel sheet having a tensile strength of 700 MPa or more.

In general, in order to increase the rigidity of the part, it is effective to change the welding conditions such as changing the shape of the part or increasing the number of welding points or converting the laser welding or the like to a part capable of spot welding. However, when these parts are used in automobiles, it is not easy to change the shape of the components in the limited space inside the vehicle, and changing the welding conditions causes problems such as an increase in cost.

As a result, it is effective to increase the Young's modulus of the material used for the part in order to increase the rigidity of the part without changing the shape of the part or the welding conditions.

In general, the stiffness of a part with the same shape and welding conditions of the part is expressed by the Young's modulus of the material and the geometric moment of inertia of the part. In addition, the geometric moment of inertia can be expressed as approximately proportional to t ^λ when the thickness of the material is t. (Lambda) is thickness sensitivity and has a value of 1-3 according to the shape of a component. For example, for one plate shape, such as a panel part of an automobile, λ is about 3, while for a column shape like a structural part, λ is about 1.

If the λ of the part is 3, if you want to reduce the thickness by 10% while maintaining the stiffness of the part, you must improve the Young's modulus of the material by 37%. An increase of 11% is sufficient.

That is, in the case of a part having a lambda close to 1 such as a columnar part, it is very effective to increase the Young's modulus of the steel sheet itself to reduce the weight. In particular, in the high strength and thin steel sheet, it is very demanded to greatly increase the Young's modulus of the steel sheet.

In general, the Young's modulus is greatly influenced by the atomic structure, and is known to be the highest in the closest direction of the atom. Therefore, in the steel process including rolling and heat treatment by rolls, it is effective to develop the {112} <110> direction in order to develop a favorable generation of the Young's modulus of the steel, which is a centered cubic lattice. It can be increased in the vertical direction.

Steel plates have been tested in various ways by controlling the atomic structure to increase the Young's modulus.

For example, Patent Document 1 discloses that steel in which Nb or Ti is added to ultra-low carbon steel is hot rolled at a rolling ratio of 85% or higher at Ar ₃ -Ar ₃ + 150 ° C, thereby promoting the transformation of ferrite in microcrystalline austenite. In the hot rolled steel sheet step, the ferrite structures are {311} <011> and {332} <113>, which are cold rolled and recrystallized annealed in the initial direction, and the rolling direction is set to {211} <011>. A technique for increasing the Young's modulus in a direction perpendicular to is disclosed.

Further, in Patent Document 2, the C content is 0.02 - in the rolling rate at 950 ℃ over 50% - Ar ₃ to the addition of Nb, Mo, B to 0.15% of low-carbon steel, and developing the [211] <011> A method for producing a hot rolled steel sheet for increasing Young's modulus is disclosed.

Furthermore, the rolling reduction ratio of under Ar ₃ transformation point in the Patent Document 3, addition of Si and Al in the low carbon steel of a C content of 0.05 or less to enhance Ar ₃ transformation point, and hot rolling to 60% or more, the rolling direction A method for producing a hot rolled steel sheet for increasing the Young's modulus in the direction perpendicular to is disclosed.

Patent Document 1: Japanese Patent Application Laid-Open No. 05-255804

Patent Document 2: Japanese Patent Application Laid-Open No. 08-311541

Patent Document 3: Japanese Patent Application Laid-Open No. 09-053118

However, the above-described techniques have the following problems.

In the technique disclosed in Patent Literature 1, the Young's modulus of the steel sheet is increased by using ultra low carbon steel having a C content of 0.01% or less to control the texture, but the tensile strength is lower than about 450 MPa. There is a problem to increase.

In the technique disclosed in Patent Document 2, since the C content is as high as 0.02-0.15%, it is possible to increase the strength, but because the target steel sheet is a hot rolled steel sheet, it becomes difficult to control the aggregate structure through cold working and further increase the Young's modulus. In addition to the problem that it is difficult to make, it is difficult to stably manufacture a high strength steel sheet having a thickness of less than 2.0 mm by low temperature finish rolling.

Further, in the technique disclosed in Patent Document 3, there is a problem that the grain boundaries are coarsened by rolling in the ferrite region, and the workability is significantly deteriorated.

Therefore, in order to increase the Young's modulus of the steel sheet in the prior art, it is intended for the thick hot rolled steel sheet or the mild steel sheet, and it is difficult to increase the Young's modulus of the high strength thin steel sheet having a thickness of 2.0 mm or less using the conventional technique.

As a reinforcing mechanism for increasing the tensile strength of the steel sheet to 590 MPa or more, there are mainly a precipitation reinforcing mechanism and a transformation structure reinforcing mechanism.

When the precipitation strengthening mechanism is used as the reinforcing mechanism, it is possible to increase the strength by making the decrease in the Young's modulus of the steel sheet as low as possible, but the following difficulties arise. That is, for example, in the case of using the precipitation reinforcing mechanism for finely precipitating carbonitrides such as Ti and Nb, in the hot rolled steel sheet, fine precipitation can be performed during winding after hot rolling to obtain an increase in strength. Coarsening of precipitates in the recrystallization annealing step after cold rolling is inevitable, and it is difficult to increase the strength through precipitation strengthening.

In the case of using the metamorphic tissue reinforcing mechanism as the reinforcing mechanism, there is a problem that the Young's modulus of the steel sheet is lowered due to the deformation induced in the low temperature transformation phase such as the bainite phase or the martensite phase.

Accordingly, an object of the present invention is to solve the above problems and provide a high strength high strength steel sheet having a tensile strength of 590 MPa or more, preferably 700 MPa or more, a Young's modulus of 230 GPa or more, preferably 240 GPa or more, and a plate thickness of 2.0 mm or less. It is to provide an advantageous method for making it.

In order to achieve the above object, the main configuration of the present invention is as follows.

(I) As mass%, C: 0.02 to 0.15%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less And Ti: 0.02 to 0.50% and the C, N, S and Ti contents are represented by the following formulas (1) and (2), that is,

Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S≥0.01 (1)

0.01 ≤ C- (12 / 47.9) x Ti * ≤0.05 (2)

And the remainder are substantially iron and unavoidable impurities, and the structure has a ferrite phase as a main phase and a martensite phase of 1% or more by area ratio, a tensile strength of 590 MPa or more and a Young's modulus of 230 GPa or more. High strength high strength steel sheet made.

(II) The compound (I) further contains one or two of Nb: 0.005 to 0.04% and V: 0.01 to 0.20% as mass%, in addition to the above composition, satisfying the formula (1). And instead of the above formula (2),

0.01 ≤ C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / 50.9) × V≤0.05 (3)

High strength high strength steel sheet, characterized in that to satisfy.

(III) In the case of (I) or (II), in addition to the above composition, Cr: 0.1% to 1.0%, Ni: 0.1% to 1.0%, Mo: 0.1% to 1.0%, and Cu: 0.1% to% by mass High strength high strength steel sheet, characterized in that it further contains at least one of 2.0% and B: 0.0005 to 0.0030%.

(IV) As mass%, C: 0.02 to 0.15%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less And Ti: 0.02 to 0.50%, wherein the C, N, S and Ti content is expressed by the following formulas (1) and (2), namely

Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S≥0.01 (1)

0.01 ≤ C- (12 / 47.9) x Ti * ≤0.05 (2)

The steel material having a composition satisfying the above characteristics is subjected to the hot rolling step under the condition that the total reduction ratio at 950 ° C. or lower is 30% or more and the finish rolling ends at 800 to 900 ° C., and the hot rolled sheet is lower than 650 ° C. After winding up and pickling, cold rolling was carried out at a reduction ratio of 50% or more, the temperature was raised from 500 ° C to 780 ° C to 900 ° C at a temperature increase rate of 1 to 30 ° C / s, and then cooled to 5 ° C / s or more. A method of producing a high rigidity high strength steel sheet, which is cooled to 500 ° C. at a speed and subjected to annealing.

(V) In the above (IV), the steel material further contains one or two of Nb: 0.005 to 0.04% and V: 0.01 to 0.20% as mass% in addition to the composition, and the formula (1) ) And satisfying the following equation (3) instead of equation (2):

0.01 ≤ C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / 50.9) × V≤0.05 (3)

Method for producing a high strength high strength steel sheet, characterized in that to satisfy.

(VI) The steel material according to (4) or (5), wherein, in addition to the composition, the steel material is Cr: 0.01 to 1.0%, Ni: 0.1 to 1.0%, Mo: 0.1 to 1.0%, Cu: 0.1-2.0% and B: 0.0005-0.0030% The manufacturing method of the high rigidity high strength steel sheet characterized by further containing 1 or more types.

According to the present invention, it is possible to provide a high rigidity high strength steel sheet having a tensile strength of 590 MPa or more and a Young's modulus of 230 GPa or more.

That is, a low carbon steel material containing Mn and Ti is rolled to promote transformation of unrecrystallized austenite into ferrite at a temperature of 950 ° C. or lower in hot rolling, and then cold to develop crystal directions useful for improving Young's modulus. After rolling, a low temperature transformation phase is generated that controls the heating rate in the annealing process and prevents the decrease of the Young's modulus in the cooling step by cracking the two phase regions, and most of the ferrite phase useful for improving the Young's modulus remains. Through this, a thin steel sheet that satisfies high strength and high Young's modulus is produced, which exhibits an industrially effective effect.

In more detail, in the hot rolling, a low-carbon steel material containing Mn and Ti is subjected to rolling shrinkage in an austenite low temperature region, thereby increasing the unrecrystallized austenite structure in the crystal direction of {112} <111>. The ferrite transformation is promoted from the unrecrystallized austenite of {112} <111> to sequentially promote the ferrite direction of {113} <110> in the cooling process.

In cold rolling after winding and pickling, rolling is carried out at a reduction ratio of 50% or more to rotate the crystallization direction of {113} <110> in the {112} <110> direction which is useful for improving the Young's modulus, and in the subsequent annealing process In the temperature rising process, the temperature is heated to the cracking temperature at a heating rate of 500 ° C. to 1-30 ° C./s to promote recrystallization of ferrite having a {112} <110> direction, and at the same time, a part of {112} <110> It is possible to reach the two phase region with the unrecrystallized grain boundary of and to promote the austenite transformation from the unrecrystallized ferrite of {112} <110>.

In addition, in the process of transforming the austenite phase in the cooling process after the cracking treatment into a ferrite phase, the ferrite grain boundary having the direction of {112} <110> is increased to enhance the Young's modulus, and the steel having high hardenability by Mn addition is increased. Cooling at rates above &lt; RTI ID = 0.0 &gt; C / s &lt; / RTI &gt; produces a low temperature transformation phase, thereby increasing strength.

In addition, the low temperature transformation phase is transformed again when the austenite phase transformed from the ferrite having the {112} <110> direction is cooled, so that the {112} <110> direction may be developed even in the crystal direction of the low temperature transformation phase.

Therefore, the Young's modulus is increased by developing {112} <110> in the ferrite phase, and in particular, the {112} <110> direction is increased in the direction of the low temperature transformation phase which greatly affects the decrease of the Young's modulus, thereby producing a low temperature transformation phase. The strength can be increased by this, and the decrease in Young's modulus due to the occurrence of low temperature transformation phase can be largely suppressed.

1 is a graph showing the effect of the total reduction rate at 950 ° C. or lower at Young's modulus.

2 is a graph showing the effect of final temperature in hot finish rolling at Young's modulus.

3 is a graph showing the influence of the winding temperature on the Young's modulus.

4 is a graph showing the effect of the reduction ratio during cold rolling at Young's modulus.

5 is a graph showing the effect of the average temperature increase rate from 500 ° C. to an cracking temperature at annealing at Young's modulus.

The high rigidity high strength steel sheet according to the present invention is a steel sheet having a tensile strength of 590 MPa or more, preferably 700 MPa or more, a Young's modulus of 230 GPa or more, preferably 240 GPa or more and a thickness of 2.0 mm or less. In addition, the object steel plate in this invention includes not only a cold rolled steel plate but also the steel plate in which surface treatment, such as a hot dip galvanizing material containing an alloying, an electrogalvanizing material, etc. was performed.

The reason for limiting the component composition of the steel sheet of the present invention will be described. In addition, although the composition unit of each element of the said steel plate is the "mass%," it will be represented simply as "%."

C: 0.02-0.15%

C is an element which stabilizes austenite, and improves the hardenability in the cooling process during annealing after cold rolling, and contributes to increasing the strength by promoting the formation of low-temperature transformation phase. In addition, C contributes to increasing the Young's modulus by promoting the deformation of the ferrite grain boundary having the {112} <110> direction after cold rolling the unrecrystallized ferrite to austenite in the temperature rising process of the annealing step.

To obtain this effect, the C content needs to be at least 0.02%, preferably at least 0.05%, more preferably at least 0.06%. On the other hand, when the C content is more than 0.15%, the fraction of hard low temperature transformation phase becomes large, the strength of the steel is extremely increased and the workability is weakened. In addition, a large amount of C content inhibits recrystallization in a direction useful for increasing the Young's modulus in the annealing step after cold rolling. In addition, a large amount of C results in weak weldability.

For this reason, the C content needs to be 0.15% or less, preferably 0.10% or less.

Si: 1.5% or less

Si raises the Ar ₃ transformation point in hot rolling, so that when the rolling is finished at 800-900 ° C, if the Si content exceeds 1.5%, it becomes difficult to roll in the austenite region and to increase the Young's modulus. You will not be able to get directions. In addition, a large amount of Si weakens the weldability of the steel sheet, but accelerates the generation of a so-called red scale surface pattern by promoting the production of fayalite on the surface of the slab during heating in the hot rolling process step. In addition, when used as a cold rolled steel sheet, the Si oxide produced on the surface weakens chemical deformation workability, and when used as a hot dip galvanized steel sheet, the Si oxide causes unplating. Therefore, the Si content needs to be 1.5% or less. In the case of a steel sheet or a hot dip galvanized steel sheet requiring surface properties, the Si content is preferably 0.5% or less.

In addition, Si is an element that stabilizes ferrite, and in the cooling process after cracking two phase regions in the annealing step after cold rolling, the transformation of ferrite is promoted to enrich C in the austenite, whereby austenite is stabilized. And can promote the formation of low temperature transformation phases. Because of this, the strength of the steel can be increased as necessary. In order to obtain such an effect, the Si content is preferably 0.2% or more.

Mn: 1.0-3.5%

Mn is one of the important elements of the present invention. Mn plays a role in suppressing the recrystallization of the processed austenite in hot rolling. In addition, Mn may promote the transformation of unrecrystallized austenite to ferrite to develop the {113} <110> direction, and may improve the Young's modulus in subsequent cold rolling and annealing steps.

In addition, Mn as a stabilizing element of austenite lowers the Ac ₁ transformation point in the temperature rising process in the annealing step after cold rolling, promotes the transformation of unrecrystallized ferrite into austenite, and the low temperature transformation phase generated in the cooling process after cracking treatment. It is possible to develop a direction useful for improving the Young's modulus so that the formation of the low temperature transformation phase with respect to the direction can be controlled at the same time.

In addition, Mn can improve the hardenability during the cracking treatment in the annealing step and the cooling process after the annealing, thereby greatly promoting the formation of the low temperature transformation phase and greatly contributing to the increase in strength. In addition, Mn can act as a solid solution strengthening element and contribute to increasing the strength of the steel. In order to obtain this effect, the content of Mn needs to be 1.0% or more, preferably 1.5% or more.

On the other hand, when the Mn content exceeds 3.5%, the Ac ₃ transformation point is greatly reduced during the temperature rising process of the annealing step after cold rolling, and as a result, it becomes difficult to recrystallize the ferrite phase in both regions and to the austenite single phase region above the Ac ₃ transformation point. It is necessary to raise the temperature. Thus, the Young's modulus is lowered without developing ferrite in the {112} <110> direction which is useful for increasing the Young's modulus obtained through recrystallization of the processed ferrite. In addition, a large amount of Mn reduces the weldability of the steel sheet. In addition, a large amount of Mn increases the deformation resistance of the steel in hot rolling, which increases the rolling load and causes difficulties in operation. Therefore, the Mn content is 3.5% or less.

P: 0.05% or less

Since P is separated from the grain boundary, when the P content is more than 0.05%, the ductility and toughness of the steel sheet is reduced, and the weldability is also weakened. In the case of using an alloy hot-dip galvanized steel sheet, the alloying speed is delayed by P. Therefore, the P content needs to be 0.05% or less. On the other hand, P is an effective element for increasing the strength as a solid solution strengthening element, and serves to promote the concentration of C in the austenite as a ferrite stabilizing element. In the steel to which Si is added, it also serves to suppress the occurrence of red scale. In order to acquire such an effect, it is preferable to make P content into 0.01% or more.

S: 0.01 or less

S lowers the hot toughness, causing hot rupture and significantly weakens the surface properties. In addition, S hardly affects the strength, but forms coarse MnS as an impurity element, thereby weakening ductility and drill-spreading properties. This problem becomes remarkable when the content of S exceeds 0.01%, so it is desirable to reduce the S content as much as possible. Therefore, the S content is 0.01% or less. In view of improving the drill spreading property, the S content is preferably 0.005% or less.

Al: 1.5% or less

Al is a useful element for deoxidized steel to improve the cleanliness of the steel. However, Al is a ferrite stabilizing element and greatly improves the Ar ₃ transformation point of the steel, and as a result, when finishing rolling at 800-900 ° C., if the Al content exceeds 1.5%, the rolling in the austenitic region is It becomes difficult to suppress the development of the crystal direction necessary for increasing the Young's modulus. Therefore, the Al content needs to be 1.5% or less. In view of this, the Al content is preferably low, and more preferably limited to 0.1% or less. On the other hand, Al, a ferrite forming element, promotes the formation of ferrite in the cooling process after cracking the two phase regions in the annealing step after cold rolling to concentrate C in the austenite. Through this, austenite is stabilized to promote the formation of low temperature transformation phase. As a result, the strength of the steel can be improved if necessary. In order to obtain such an effect, the Al content is preferably 0.2% or more.

N: 0.01% or less

N is a detrimental element because it is accompanied by slab breakage which causes surface defects during hot rolling. When the N content exceeds 0.01%, the occurrence of slab breakage and surface defects becomes remarkable. In addition, when carbonitride-forming elements such as Ti and Nb are added, N forms coarse nitride at a high temperature to suppress the effect of addition of the carbonitride-forming element. Therefore, the N content needs to be 0.01% or less.

Ti: 0.02-0.50%

Ti is the most important element in this invention. That is, Ti controls the recrystallization of austenite processed in the finish rolling process in hot rolling to promote the transformation of ferrite from unrecrystallized austenite and to develop the {113} <110> direction. In addition, Ti increases Young's modulus in subsequent cold rolling and annealing steps. In addition, the recrystallization of the ferrite processed during the elevated temperature during the annealing step after the cold rolling is suppressed to promote the transformation of austenite in the unrecrystallized ferrite, and is useful for increasing the Young's modulus with respect to the low-temperature transformation phase generated during the cooling process after the cracking treatment. Aroma is developed, and the fall of the Young's modulus by formation of a low temperature transformation phase can be suppressed. In addition, fine Ti carbide can contribute to an increase in strength. In order to obtain such an action, the Ti content needs to be 0.02% or more, preferably 0.03% or more.

On the other hand, when the Ti content is more than 0.50%, all the carbide nitrides cannot be solid solution upon reheating in the usual hot rolling step, and coarse carbide nitrides remain, which inhibits recrystallization of the austenite processed in the hot rolling step. The effect or the effect of suppressing recrystallization of the processed ferrite in the annealing step after cold rolling cannot be obtained. Further, even when the hot rolling of the continuous post-slab is started without cooling after the continuous post-slab cooling, the improvement of the recrystallization inhibiting effect is not recognized when the Ti content exceeds 0.50%. An increase in alloy cost is incurred. Therefore, the Ti content needs to be 0.50% or less, preferably 0.20% or less.

In the present invention, the contents of C, N, S and Ti should satisfy the relationship of the following formulas (1) and (2).

Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S≥0.01 (1)

0.01 ≤ C- (12 / 47.9) x Ti * ≤0.05 (2)

Ti tends to form coarse nitrides and sulfides in the high temperature region. The formation of such nitrides and sulfides results in a decrease in the effect of inhibiting recrystallization by Ti addition. Therefore, the amount of Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S, which is a Ti content not fixed with nitrides and sulfides, needs to be 0.01% or more, preferably 0.02% or more.

If an unfixed C content as carbonitride is present in excess of 0.05%, the introduction of strain during cold rolling is not constant, and recrystallization in a direction useful for increasing the Young's modulus is suppressed, resulting in (C- (12 / 47.9 The content of C which is not fixed with carbides calculated by) x Ti *) needs to be 0.05% or less. On the other hand, when the C content not fixed with carbide is less than 0.01%, the C content in the austenite decreases upon annealing in two phase regions after cold rolling, thereby increasing the strength by inhibiting the formation of martensite after cooling. It becomes difficult. Therefore, the amount of C- (12 / 47.9) x Ti *, which is not fixed as carbide, is 0.01-0.05% or more.

In addition, the expression "residual substantially iron and unavoidable impurities" used herein means that the inclusion of other trace elements is included in the scope of the present invention without impairing the operation and effects of the present invention. In order to further increase the strength, in addition to the above limitation of the chemical composition, if necessary, one or two of Nb: 0.005-0.04% and V: 0.01-0.20%, but Cr, Ni, Mo, Cu and B You may add 1 or more types of components chosen from.

Nb: 0.005 to 0.04%

Nb is an element which contributes to the increase in strength by forming fine carbide nitride. It is also an element that promotes ferrite transformation from unrecrystallized austenite and increases Young's modulus by suppressing recrystallization of austenite processed in the finish rolling step in hot rolling. In order to obtain such an effect, it is preferable to make Nb content 0.005% or more. On the other hand, when the Nb content exceeds 0.04%, the rolling load in hot rolling and cold rolling greatly increases and manufacturing difficulties occur, so that the Nb content is 0.04% or less, more preferably 0.01% or less. It is preferable to set it as.

V: 0.01-0.20%

V is an element which contributes to the increase in strength by forming fine carbide nitride. In order to have such an effect, it is preferable to make content of V into 0.01% or more. On the other hand, when the V content exceeds 0.20%, the effect of increasing the strength by more than 0.20% is small, resulting in an increase in the alloy cost.

Therefore, the content of V is preferably set to 0.01-0.20%.

In the present invention, when Ti contains Nb and / or V, the content of C, N, S, Ti, Nb and V needs to satisfy the relational formula shown in the following formula (3) instead of the above formula (2). .

0.01 ≤ C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / 50.9) × V≤0.05 (3)

Nb and V form carbides to form carbides that reduce the amount of C not fixed as carbide. Therefore, in order to make the C content not fixed as carbide to 0.01-0.05%, when Nb and / or V are added, C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / The value of 50.9) xV needs to be 0.01-0.05%.

Cr: O.1-1.O%

Cr is an element that suppresses the formation of cementite and enhances the hardenability, and can greatly contribute to increasing the strength by greatly promoting the formation of the low-temperature transformation phase in the cooling process after the cracking treatment in the annealing step. In addition, recrystallization of the austenite processed in the hot rolling step is suppressed to promote ferrite transformation from unrecrystallized austenite to develop the {113} <110> direction, and the Young's modulus may be increased in the subsequent cold rolling and annealing step. have. In order to acquire such an effect, it is preferable to contain Cr 0.1% or more. On the other hand, when the Cr content is more than 1.0%, the above effect is saturated, and the alloy cost is increased, so that Cr is preferably contained at less than 1.0%. In addition, when using the thin steel plate of this invention as a hot-dip galvanized steel plate, since Cr oxide produced on the surface causes non-plating, it is preferable to contain Cr below 0.5%.

Ni: 0.1-1.0%

Ni is an element which stabilizes austenite to increase the hardenability, and can greatly contribute to the increase in strength by greatly promoting the formation of the low temperature transformation phase in the cooling process after the cracking treatment in the annealing step. In addition, Ni, an austenite stabilizing element, lowers the Ac ₁ transformation point in the temperature rising process during the annealing step after cold rolling, promotes austenite transformation from unrecrystallized ferrite, and corresponds to the low temperature transformation phase generated during the cooling process after the cracking treatment. Therefore, the direction favorable to the increase of Young's modulus is developed. As a result, a decrease in Young's modulus can be suppressed as the low temperature transformation phase is generated. In addition, Ni suppresses recrystallization of austenite processed during hot rolling to promote ferrite transformation from unrecrystallized austenite to develop a {113} <110> direction, thereby increasing the Young's modulus in subsequent cold rolling and annealing steps. Can be. In the case of Cu-added steel, surface defects are caused by cracks due to a decrease in thermal ductility during hot rolling, but generation of surface defects can be controlled by adding Ni. In order to acquire such an effect, it is preferable that Ni is contained 0.1% or more.

On the other hand, when the content of Ni exceeds 1.0%, the Ac ₃ transformation point extremely decreases in the temperature rising process of the annealing step after cold rolling, making it difficult to recrystallize the ferrite phase in the two phase region, and the austenitic single layer having the Ac ₃ transformation point or more. It is necessary to raise the temperature to the phase region. As a result, the ferrite in a direction useful for increasing the Young's modulus obtained by recrystallization of the processed ferrite is not developed, resulting in a decrease in Young's modulus. In addition, the cost of the alloy also increases. Therefore, Ni is preferably contained at 1.0% or less.

Mo: 0.1-1.0%

Mo is an element which improves hardenability by making the mobility of an interface small, and can greatly contribute to increasing strength by greatly promoting formation of a low-temperature transformation phase in the cooling process of the annealing step after cold rolling. In addition, recrystallization of the processed austenite can be suppressed, and the transformation of ferrite from unrecrystallized austenite can be promoted to develop the {113} <110> direction, and the Young's modulus can be increased in subsequent cold rolling and annealing steps. have. In order to obtain such an action, it is preferable that Mo is contained in 0.1% or more. On the other hand, when the Mo content exceeds 1.0%, since the effect is saturated and the alloy cost increases, Mo is preferably contained at 1.0% or less.

B: 0.0005-0.0030%

B is an element which suppresses the transformation from the austenite phase to the ferrite phase to enhance the hardenability, and greatly promotes the formation of the low temperature transformation phase in the cooling process of the annealing step after cold rolling, and can greatly contribute to the increase in strength. In addition, recrystallization of the processed austenite can be suppressed and the ferrite transformation can be promoted from the unrecrystallized austenite to develop the {113} <110> direction, thereby increasing the Young's modulus in the subsequent cold rolling and annealing steps. In order to acquire such an effect, it is preferable that B is contained in 0.0005% or more. On the other hand, when the B content is more than 0.0030%, since the deformation resistance during hot rolling increases, the rolling load is increased and operation difficulties occur, so that B is preferably contained at 0.0030% or less.

Cu: 0.1-2.0%

Cu is an element which improves hardenability and can greatly contribute to increasing strength by greatly promoting the formation of low-temperature transformation phase in the cooling process of the annealing step after cold rolling. In order to acquire such an effect, it is preferable to contain Cu 0.1% or more. On the other hand, when the Cu content exceeds 2.0%, thermal ductility is lowered, surface defects due to cracks during hot rolling are caused, and the curing effect by Cu is saturated, and therefore Cu is preferably contained at 2.0% or less. .

The reason for limitation of the tissue according to the present invention will be described below.

In the thin steel sheet of the present invention, it is required to have a structure containing a ferrite phase as a main phase and having a martensite phase of 1% or more in area ratio.

As used herein, the term 'ferrite phase as the main phase' means that the area ratio of the ferrite phase is 50% or more.

Since the ferrite phase is less deformed, is useful for increasing the Young's modulus, is excellent in ductility, and has good workability, the structure is required to be a ferrite phase as a main phase.

Moreover, in order to make tensile strength of a steel plate 590 Mpa or more, it is necessary to form the low temperature transformation phase which is a hard phase in the part other than the ferrite phase which is a main phase, or what is called a 2nd phase, and provide a composite phase. Here, the characteristic that the hardened martensite phase between the low temperature transformation phases is particularly present inside the structure is advantageous because the fraction of the second phase for obtaining the target tensile strength level is made small and the fraction of the ferrite phase is made large. Through this, it is possible to increase the Young's modulus and the workability can be improved. For this reason, the martensite phase needs to be 1% or more in area ratio with respect to the whole tissue. In order to obtain the strength of 700 MPa or more, it is preferable that the area ratio of the martensite phase is 16% or more.

The structure of the steel sheet according to the present invention is preferably a structure including a ferrite phase and martensite phase, but the ferrite phase and martensite such as bainite phase, residual austenite phase, pearlite phase, or cementite phase Even if phases other than a site phase exist in an area ratio of 10% or less, preferably 5% or less, there is no problem. That is, the sum of the area ratios of the ferrite phase and the martensite phase is preferably 90% or more, more preferably 95% or more.

Next, the reason for the limited manufacturing conditions and preferable manufacturing conditions for obtaining the high rigidity high strength steel sheet which concerns on this invention are demonstrated.

Since the composition of the steel raw material used for the manufacturing method of this invention is the same as that of the steel plate mentioned above, description of the reason for limitation of a steel raw material composition was abbreviate | omitted.

The steel sheet according to the present invention is a hot rolling step of hot rolling to a steel material having the same composition as the steel sheet composition as a hot rolled sheet, and cold rolling after pickling the hot rolled sheet to form a cold rolled sheet and The cold rolled sheet may be manufactured by annealing in which recrystallization and complex organization are performed.

(Hot rolling step)

Finish rolling: The total rolling reduction is made 30% or more at 950 ° C or less, and the rolling is finished at 800 to 900 ° C.

In the finish rolling of the hot rolling step, rolling is carried out at low temperature to develop unrecrystallized austenite structure having a crystallographic direction of {112} <111>, and in the subsequent cooling process, {112} <111> unrecrystallized austenite Is transformed into ferrite to develop the {113} <110> ferrite direction. This direction advantageously acts to improve the Young's modulus in the formation of the tissue in subsequent cold rolling and annealing steps. In order to obtain such an effect, the total pressure drop at 950 ° C. or lower is 30% or more, and finish rolling must be finished at 900 ° C. or lower. On the other hand, when the finishing temperature of finish rolling is lower than 800 degreeC, rolling load will increase remarkably by the increase of deformation resistance, and manufacturing difficulty will follow. Therefore, the finishing temperature of finish rolling needs to be 800 degreeC or more.

Winding temperature: below 650 ℃

When the coiling temperature after finish rolling exceeds 650 ° C, the carbon carbide of Ti is coarsened, and the effect of inhibiting recrystallization of ferrite becomes small during the temperature rising process of the annealing step after cold rolling, and the transformation from unrecrystallized ferrite to austenite is performed. It becomes difficult. As a result, it becomes impossible to control the direction of the low temperature transformation phase transformed in the cooling process after the cracking treatment, and the Young's modulus is greatly reduced by the low temperature transformation phase with such deformation. Therefore, the winding temperature after finishing rolling needs to be 650 degrees C or less. In addition, when the said winding temperature is too low, since many hard low temperature transformation phases generate | occur | produce and a load increases in subsequent cold rolling, it is preferable to set it as 400 degreeC or more.

(Cold rolling process)

After pickling, cold rolling is carried out at a rolling reduction of at least 50%.

After the hot rolling step, pickling is carried out to remove the scale formed on the surface of the steel sheet. Pickling is done according to conventional methods. Thereafter, cold rolling is performed. Here, by performing cold rolling at a rolling reduction rate of 50% or more, the {113} <110> direction developed in the hot rolled steel sheet may be rotated in the {112} <110> direction effective for increasing the Young's modulus. Therefore, as the {112} <110> direction is developed by cold rolling, the {112} <110> direction of the ferrite is strengthened in the tissue after the subsequent annealing step, and also in the {112} <110> direction on the low temperature transformation. Is developed. This can increase the Young's modulus. In order to acquire such an effect, it is necessary to make the reduction ratio at the time of cold rolling into 50% or more.

(Anneal phase)

Heating rate from 500 ° C to cracking temperature: 1 to 30 ° C / s, cracking temperature: 780 to 900 ° C

The rate of temperature increase in the annealing step is an important treatment condition in the present invention. In the annealing step, in the process of raising the temperature of the two-phase cracking temperature or the cracking temperature 780 to 900 ° C, recrystallization of the ferrite having the {112} <110> direction is promoted and the ferrite having the {112} <110> direction Some of the particles reach the head-phase region in the state of unrecrystallization. This facilitates the transformation of the unrecrystallized ferrite having the {112} <110> direction. Therefore, when the austenite is transformed into ferrite during cooling after the cracking treatment, the Young's modulus can be increased by promoting the growth of the particles of the ferrite having the {112} <110> direction. In addition, when the strength is increased by generating a low temperature transformation phase, the austenite phase transformed from the ferrite including the {112} <110> direction is re-transformed upon cooling, so that the {112} <110> direction is determined as the low temperature transformation phase. Can be developed with respect to the direction. By developing the {112} <110> direction of the ferrite phase, the Young's modulus is increased, and in particular, the {112} <110> direction is developed in the low temperature transformation phase direction which greatly affects the decrease of the Young's modulus. As a result, a decrease in Young's modulus due to the formation of the low temperature transformation phase can be suppressed while the low temperature transformation phase is generated. In the case of austenite transformation from unrecrystallized ferrite while promoting ferrite recrystallization during heating, the average temperature increase rate from 500 ° C. to 780-900 ° C., which is a crack treatment temperature, greatly affects the recrystallization behavior. Must be / s. The reason why the cracking treatment temperature is 780 to 900 ° C. is that unrecrystallized structure remains when it is lower than 780 ° C., and when it exceeds 900 ° C., the amount of austenite produced increases, which is advantageous for increasing the Young's modulus. This is because it becomes difficult to develop ferrite having a> direction.

In addition, the cracking treatment time is not particularly limited, but it is preferable to set it to 30 seconds or more in producing austenite, and if it is too long, the production efficiency becomes worse, so it is preferable to set it to about 300 seconds or less.

Cooling rate up to 500 ℃ after cracking: 5 ℃ / s or more

It is necessary to create a low temperature transformation phase including the martensite phase in order to increase the strength in the cooling process after the crack treatment. For this reason, it is necessary to make the average cooling rate to 500 degrees C / s or more after a cracking process.

In the present invention, the steel having the chemical component according to the target strength level is first dissolved. As the melting method, a conventional converter method, an electric furnace method, or the like may be appropriately applied. The molten steel is cast into slabs, followed by hot rolling as is or by cooling and heating. After finishing by the above-described finishing conditions in hot rolling, the steel sheet is wound at the winding temperature described above, followed by normal pickling and cold rolling. Regarding the annealing, the cooling rate can be increased to a range in which the target low temperature transformation phase can be obtained upon cooling after the cracking treatment by raising the temperature under the above-described conditions. Thereafter, in the case of a cold rolled sheet, aging treatment may be performed, or when manufactured as a hot-dip galvanized steel sheet, it may be passed through molten zinc, and when manufactured as an alloyed hot-dip galvanized steel sheet, 500 ° C for alloying treatment. Reheating may be performed up to the above temperature.

Example

An embodiment of the present invention will be described. This invention is not limited to these Examples.

First, steel A having the chemical composition shown in Table 1 was dissolved in a laboratory vacuum melting furnace, and cooled to room temperature once to produce a steel mass (steel material).

Table 1

Thereafter, hot rolling, pickling, cold rolling, and annealing were sequentially performed in a laboratory. Basic manufacturing conditions are as follows. The steel mass was heated at 1250 ° C. for 1 hour and then hot rolled to obtain a total pressure reduction rate of 950 ° C. or lower at 40%, and a final rolling temperature (corresponding to the final temperature of finish rolling) at 860 ° C. A hot rolled sheet of 4.0 mm was obtained. Thereafter, the hot rolled sheet was held at a maximum of 600 ° C. to be subjected to winding conditions (corresponding to 600 ° C. winding temperature), put in a furnace at 600 ° C. for 1 hour, and then cooled in the furnace. The hot rolled sheet thus obtained was pickled, cold rolled at a reduction ratio of 60% to a plate thickness of 1.6 mm, and then heated to 500 ° C. at an average of 10 ° C./s, and further at 500 ° C. at an average of 5 ° C. It heated up to the cracking temperature of 820 degreeC by s. Next, after the cracking treatment at 820 ℃ for 180 seconds, cooled to 500 ℃ at an average cooling rate of 10 ℃ / s and maintained at 500 ℃ for 80 seconds, the steel sheet was air-cooled.

In this experiment, the manufacturing conditions which are basic conditions were changed individually to the following conditions. That is, 20 to 60% of the total pressure reduction rate at 950 degreeC or less, 800-920 degreeC of the final temperature of hot finishing rolling, 500-670 degreeC of the coiling temperature, 40-75% of the rolling reduction of annealing, The average temperature increase rate from 500 degreeC to the cracking process temperature (820 degreeC) was made into 0.5-35 degreeC / s, and it experimented on the basic conditions except the individual condition which changed.

From the sample after annealing, a 10 mm x 120 mm specimen was cut out in a longitudinal direction perpendicular to the longitudinal rolling direction, and finished to a thickness of 0.8 mm by mechanical polishing and chemical polishing to remove deformation. The resonant frequency of the specimen was measured using a lateral oscillating internal friction measuring device and the Young's modulus was calculated therefrom. Tensile test of JIS 5 was cut | disconnected in the direction perpendicular | vertical to the rolling direction with respect to the board which carried out the 0.5% temper rolling, and the tension test was done. In addition, the cross-sectional structure was observed by scanning electron microscopy (SEM) after corroding with nital, and after taking three photographs in a viewing area of 30 µm x 30 µm, ferrite phase and The area ratio of martensite phase was measured and the average value for each phase was calculated | required as the area ratio (fraction) of each phase.

  As a result, the mechanical property values under the basic conditions in the experiment according to the production method of the present invention are Young's modulus E: 242GPa, TS: 780MPa, El: 23% and ferrite phase fraction: 67%, martensite phase fraction: 28% The steel sheet was found to have an excellent strength-ductility balance and excellent Young's modulus.

In the above structure, the remainder other than the ferrite phase and the martensite phase is any of bainite phase, residual austenite phase, pearlite phase and cementite phase.

Hereinafter, the relationship between a manufacturing condition and a Young's modulus is demonstrated based on a test result with reference to drawings. Here, under any of the experimental conditions, the tensile strength was 730 to 820 MPa, the ferrite phase fraction was 55 to 80%, the martensite phase fraction was 17 to 38%, and the balance was bainite phase, residual austenite phase, pearlite phase, and cementene. One of the tightest statues.

In FIG. 1, the influence of the total pressure reduction rate in 950 degreeC or less on Young's modulus is shown. When the total reduction ratio was 30% or more, which is the preferred range of the present invention, the Young's modulus showed an excellent value of 230 GPa or more.

In FIG. 2, the influence of the final temperature of hot finish rolling is shown on a Young's modulus. In the case where the final temperature is 900 ° C. or lower, which is a preferred range of the present invention, the Young's modulus showed an excellent value of 230 GPa or more.

In FIG. 3, the influence of the winding temperature on the Young's modulus is shown. When the coiling temperature was below 650 ° C., which is a preferred range of the present invention, the Young's modulus showed an excellent value of 230 GPa or more.

In FIG. 4, the influence of the reduction ratio in cold rolling to Young's modulus is shown. When the reduction ratio is 50% or more, which is a preferred range of the present invention, the Young's modulus showed an excellent value of 230 GPa or more.

In FIG. 5, the influence of the average temperature increase rate from 500 degreeC at the time of annealing to 820 degreeC which is a cracking process temperature at an Young's modulus is shown. When the temperature increase rate was 1 to 30 ° C / s, which is a preferable range of the present invention, the Young's modulus showed an excellent value of 230 GPa or more.

In addition, steel B-Z and AA-AI of the chemical composition shown in Table 2 were melt | dissolved in the laboratory vacuum melting furnace, and cooled to room temperature, and the steel mass (steel material) was produced. Thereafter, hot rolling, pickling, cold rolling, and annealing were performed sequentially under the conditions shown in Table 3. The steel mass was heated at 1250 ° C. for 1 hour, then hot rolled to roll at various rolling temperatures to obtain a hot rolled sheet having a plate thickness of 4.0 mm. Thereafter, after the desired coiling temperature was reached, the resultant was put in a furnace, held for 1 hour, the furnace was cooled, and winding conditions were performed. The hot rolled sheet was pickled, cold rolled at various reduction ratios, and had a plate thickness of 0.8 to 1.6 mm, and then heated up to 500 ° C. at an average of 10 ° C./s, followed by various averages shown in Table 3. The temperature was raised to the desired cracking temperature at an elevated rate. Next, after the cracking treatment was performed for 180 seconds at the cracking treatment temperature, cooling was performed at various average cooling rates shown in Table 3, held at 500 ° C for 80 seconds, and then air-cooled to room temperature.

Table 4 summarizes the characteristics obtained by the above-mentioned irradiation. In this table, the structure other than the martensite phase and the ferrite phase is any one of a bainite phase, a retained austenite phase, a pearlite phase and a cementite phase.

TABLE 2

Table 3

Table 4

For steel D, the C content is as small as 0.01%, the fraction of martensite is 0%, and TS is smaller than the preferred range of the present invention. In steel E, the content (SC) of C which is not fixed as carbide is as large as 0.08%, the fraction of ferrite phase is made as small as 30%, and the Young's modulus is smaller than the preferred range of the present invention. In the steel F, SC is as large as 0.06% and the Young's modulus is smaller than the preferred range of the present invention. In steel K, the Mn content is as high as 3.6% and the Young's modulus is smaller than the preferred range of the present invention. In steel AD, the C content is as large as 0.16%, the SC is as large as 0.14%, the ferrite phase fraction is as small as 25%, and the Young's modulus is smaller than the preferred range of the present invention. In the steel AF, the Mn content is as small as 0.9%, and the TS and Young's modulus are smaller than the preferred range of the present invention. In the steel AI, the Ti content is as small as 0.01%, the Ti * is as small as 0.00%, and the Young's modulus is smaller than the preferred range of the present invention.

In all other steels, TS and Young's modulus satisfy | filled the preferable range of this invention in all within the preferable range of this invention.

The present invention can provide a high strength high rigidity steel sheet having a tensile strength of 590 MPa or more and a Young's modulus of 230 GPa or more.

Claims (6)

As mass%, C: 0.02 to 0.15%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less and Ti: Relations containing 0.02 to 0.50% and the C, N, S and Ti contents are represented by the following formulas (1) and (2): Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S≥0.01 (1) 0.01 ≤ C- (12 / 47.9) x Ti * ≤0.05 (2) And the remainder are substantially iron and unavoidable impurities, and the structure has a ferrite phase as a main phase and a martensite phase of 1% or more by area ratio, a tensile strength of 590 MPa or more and a Young's modulus of 230 GPa or more. High strength high strength steel sheet made. The method according to claim 1, further comprising one or two of Nb: 0.005 to 0.04% and V: 0.01 to 0.20% as mass%, in addition to the composition, satisfying the formula (1) and Instead of (2) the relation of 0.01 ≤ C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / 50.9) × V≤0.05 (3) High strength high strength steel sheet, characterized in that to satisfy. The method according to claim 1 or 2, wherein, in addition to the composition, Cr: 0.1% to 1.0%, Ni: 0.1% to 1.0%, Mo: 0.1% to 1.0%, Cu: 0.1% to 2.0%, and B: A high-strength high strength steel sheet further comprising at least one of 0.0005 to 0.0030%. As mass%, C: 0.02 to 0.15%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less and Ti: 0.02 to 0.50%, and the C, N, S and Ti content is represented by the following formulas (1) and (2), that is, Ti * = Ti- (47.9 / 14) × N- (47.9 / 32.1) × S≥0.01 (1) 0.01 ≤ C- (12 / 47.9) x Ti * ≤0.05 (2) The steel material having a composition satisfying the above characteristics is subjected to the hot rolling step under the condition that the total reduction ratio at 950 ° C. or lower is 30% or more and the finish rolling ends at 800 to 900 ° C., and the hot rolled sheet is lower than 650 ° C. After winding up and pickling, cold rolling was carried out at a reduction ratio of 50% or more, the temperature was raised from 500 ° C to 780 ° C to 900 ° C at a temperature increase rate of 1 to 30 ° C / s, and then cooled to 5 ° C / s or more. A method of producing a high rigidity high strength steel sheet, which is cooled to 500 ° C. at a speed to perform annealing. The steel material according to claim 4, further comprising one or two of Nb: 0.005 to 0.04% and V: 0.01 to 0.20% as mass% in addition to the composition, and satisfies the formula (1). And instead of the above formula (2), 0.01 ≤ C- (12 / 47.9) × Ti *-(12 / 92.9) × Nb- (12 / 50.9) × V≤0.05 (3) Method for producing a high strength high strength steel sheet, characterized in that to satisfy. 6. The steel material according to claim 4 or 5, wherein, in addition to the composition, the steel is made of Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, Mo: 0.1 to 1.0%, and Cu: 0.1 to 2.0. % And B: 0.0005 to 0.0030% of the method for producing a high strength high strength steel sheet, characterized in that it further contains one or more.

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