JP4736617B2 - High-strength, high-strength cold-rolled steel sheet and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet having a tensile strength TS of 590 MPa or more used for a columnar structural member such as a center pillar of an automobile, in particular, a high-rigidity high-strength cold-rolled steel sheet and a method for producing the same.

  In recent years, in response to increasing interest in global environmental issues, there is a demand for exhaust gas regulations for automobiles, and weight reduction of automobile bodies has become an extremely important issue. In order to reduce the weight of an automobile body, it is effective to increase the strength of a steel plate used for the vehicle body and reduce the thickness of the steel plate. For this reason, recently, high-strength cold rolling with a TS of 590 MPa or more and a plate thickness of less than 2.0 mm is also used for columnar structural members such as automobile center pillars, rockers, side frames, and cross members. There is a movement to actively apply steel plates.

  Up to now, as high-strength cold-rolled steel sheet having TS of 590 MPa or more, for example, in mass%, C: 0.04 to 0.14%, Si: 0.4 to 2.2%, Mn: 1.2 to 2.4%, P: 0.02% or less, S : 0.01% or less, Al: 0.002 to 0.5%, Ti: 0.005 to 0.1%, N: 0.006% or less, further containing one or more of Nb, Mo, V in a total amount of 0.005 to 0.1%, Ti / S ≧ 5 and a slab having a composition consisting of Fe and inevitable impurities in the balance is hot-rolled at a finishing temperature (900 + 50 × Si) or higher, and after cold rolling at a reduction rate of 50 to 85%, 700 to 900 Annealed for 10 s to 5 min in the two-phase coexistence temperature range of ferrite and austenite at ℃, cooled to 250 to 500 ℃ with an average cooling rate between 700 ℃ and 500 ℃ to 1 to 120 ℃ / s, as needed After reheating, hold in the temperature range of 250 to 600 ° C for 30 s to 10 min and then cool to make the total volume ratio of martensite and austenite 6% or more, and the hard phase of martensite, residual austenite and bainite Low yield ratio high-strength cold-rolled steel sheet with excellent hole expansibility, in which the volume ratio α% of α ≦ 50000 × (Ti / 48 + Nb / 93 + Mo / 96 + V / 51) is disclosed (Patent Document 1).

Also, by weight%, C: 0.02 to 0.30%, Si: 1.50% or less, Mn: 0.60 to 3.0%, P: 0.20% or less, S: 0.05% or less, Al: 0.01 to 0.10%, Mo: 0.01 to 1.0%, Cr: 0.10 to 1.5%, Nb: 0.01 to 0.05%, Ti: 0.01 to 0.5%, B: 0.0005 to 0.0030%, Ca: containing at least one of 0.01% or less, with the balance being Fe and Steel with an inevitable impurity composition is hot-rolled at the Ac ₃ transformation point or higher, pickled and cold-rolled, and then heated and maintained above the recrystallization temperature and above the Ac ₁ transformation point in a continuous annealing hot-dip galvanizing line After that, in the period up to the molten zinc bath, it is rapidly cooled below the Ms point to generate martensite as a whole or part of the steel sheet, and then at least the molten zinc bath temperature and alloying at a temperature above the Ms point. A high-strength alloyed hot-dip galvanized steel sheet having excellent stretch flangeability, which is heated to a furnace temperature to generate tempered martensite in whole or in part, is disclosed ( Patent Document 2).

Patent Document 3 below relates to the Young's modulus of steel described in [Means for Solving the Problems] described later.
Japanese Unexamined Patent Publication No. 2002-69574 JP-A-6-93340 Japanese Patent Laid-Open No. 5-255804

  However, when the high-strength cold-rolled steel sheet described in Patent Document 1 or Patent Document 2 is actually applied to such a column-shaped structural member, sufficient rigidity cannot be secured, that is, problems such as bending or twisting occur. .

  An object of the present invention is to provide a high-strength cold-rolled steel sheet having high rigidity and TS of 590 MPa or more and a method for producing the same.

  In general, increasing the Young's modulus is effective for increasing the rigidity of the steel sheet. In addition, the Young's modulus of steel greatly depends on the texture, and in the case of plain steel with a body-centered cubic lattice, the atomic density is high in the <111> direction, and conversely, the atomic density is small <100> Since the direction is low, if the {112} <110> orientation is developed, the Young's modulus in the direction perpendicular to the rolling direction of the steel sheet can be increased. In order to obtain a cold rolled steel sheet with {112} <110> orientation, it is effective to develop the {113} <110> orientation after hot rolling (for example, Patent Document 3).

Therefore, the present inventors conducted various studies to improve the rigidity of the high-strength cold-rolled steel sheet reinforced with the low-temperature transformation phase from the viewpoint of formability, and developed the above {112} <110> orientation. It was found that it is effective to increase the Young's modulus and further set X in the following formula (1) to 230 GPa or more;
X = 0.005 × TS / (ε2-ε1) [GPa] (1)
Here, TS is the tensile strength [MPa] of the steel sheet, ε1 is the nominal strain [%] when the nominal stress is 0.40 × TS, and ε2 is the nominal strain [%] when the nominal stress is 0.45 × TS.

The present invention has been made on the basis of such knowledge, in mass%, C: 0.02 to 0.20%, 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, Nb: 0.01 to 0.40%, balance Fe and inevitable impurities, satisfying [Nb]-(93/14) × [N] ≧ 0.005, 200 to It has a microstructure composed of a ferrite phase and a low-temperature transformation phase obtained by tempering by staying in a temperature range of 500 ° C. for 60 s or more, a volume ratio of the low-temperature transformation phase is 1 to 60%, and a tensile strength is 590 MPa or more. In addition, a high-strength, high-strength cold-rolled steel sheet having X of 230 GPa or more in the above formula (1) is provided. Here, [M] represents the content [% by mass] of the element M.

  In the high-strength cold-rolled steel sheet of the present invention, Nb is preferably 0.03 to 0.40%.

  The high-strength cold-rolled steel sheet of the present invention can further contain at least one element of V: 0.01 to 0.5% and Ti: 0.01 to 0.2% by mass%.

  Further, the high-strength cold-rolled steel sheet of the present invention is further in mass%, Mo: 0.1 to 1.0%, Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, B: 0.0005 to 0.0020%. It can also contain an element or Cu: 0.1-2.0%.

The high-strength cold-rolled steel sheet of the present invention is, for example, a steel slab having the above composition, with a total reduction amount of 30% or more at 950 ° C. or less, and hot rolling at a finishing temperature of Ar ₃ transformation point to 900 ° C. Then, after winding at 650 ° C. or less and cold rolling at a reduction rate of 40% or more, (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn] ]) Heated at 1 to 600 s for 1 to 600 s, then cooled and retained in the temperature range of 700 to 500 ° C. for 20 s or more and retained in the temperature range of 200 to 500 ° C. for 60 s or tempered, or 700 to 500 ° C. It can be manufactured by a method in which it is cooled to 200 ° C. or lower after being retained in the temperature range, reheated and retained in the temperature range of 200 to 500 ° C. for 60 seconds or more and annealed under the conditions of tempering. Here, [M] represents the content [% by mass] of the element M.

  The present invention makes it possible to produce a high-strength cold-rolled steel sheet having a TS of 590 MPa or more that can ensure rigidity even when applied to a columnar structural member such as a center pillar of an automobile.

  Below, the high-strength cold-rolled steel sheet and its manufacturing method which are this invention are demonstrated in detail.

1) ingredients
Since C: C is an austenite stabilizing element, it enhances hardenability in the cooling process during annealing after cold rolling, promotes the formation of a low-temperature transformation phase, and greatly contributes to high strength. In addition, since the Ar ₃ transformation point is lowered, it enables hot rolling in a lower temperature range, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and {112} } Develop <110> orientation to improve Young's modulus. In order to obtain such an effect, the amount needs to be 0.02% or more. On the other hand, if the amount exceeds 0.20%, the hard low-temperature transformation phase increases and the formability deteriorates, and the ferrite phase decreases, so the Young's modulus decreases. Therefore, the C content is 0.02 to 0.20%, preferably 0.02 to 0.10%.

Si: Si significantly raises the Ar ₃ transformation point and promotes recrystallization of processed austenite during hot rolling, so if it contains more than 1.5%, it develops the crystal orientation necessary for high Young's modulus In addition, it deteriorates the weldability, chemical conversion properties and plating properties of the steel sheet, and promotes the occurrence of surface defects called red scale during hot rolling. Therefore, the Si content is 1.5% or less. In particular, when the steel sheet of the present invention requires good surface properties, or when hot dip galvanizing is performed on the steel sheet of the present invention, the Si content is preferably 0.5% or less. Si is a ferrite stabilizing element and has the effect of promoting ferrite transformation in the cooling process during annealing, concentrating C in austenite to stabilize austenite, and promoting the formation of a low temperature transformation phase. In order to sufficiently obtain this effect, it is desirable that the Si content be 0.1% or more.

Mn: Mn is one of the important elements in the present invention, and it suppresses recrystallization of processed austenite during hot rolling and lowers the Ar ₃ transformation point, enabling hot rolling at lower temperatures. And promote the ferrite transformation from unrecrystallized austenite to develop the {113} <110> orientation and improve the Young's modulus after annealing. Mn is also an austenite stabilizing element, so it lowers the Ac ₁ transformation point during the temperature rise process during annealing, promotes austenite transformation from unrecrystallized ferrite, and increases the Young's modulus in the subsequent cooling process. Develop advantageous ferrite orientation. Furthermore, Mn is a solid solution strengthening element, and enhances hardenability in the cooling process during annealing and promotes the formation of a low-temperature transformation phase, thus greatly contributing to an increase in strength. In order to obtain such an effect, the amount needs to be 1.0% or more. On the other hand, if the amount exceeds 3.5%, the formation of ferrite necessary for high Young's modulus during the cooling process during annealing is remarkably suppressed, and the low-temperature transformation phase increases, resulting in extremely high strength steel. As well as deterioration of weldability, weldability also deteriorates. Therefore, the Mn content is 1.0 to 3.5%, preferably 1.5 to 2.5%.

  When P: P is contained in an amount exceeding 0.05%, it segregates at the grain boundaries and lowers the ductility and toughness of the steel sheet and degrades the weldability. Moreover, when alloying hot dip galvanizing to the steel plate of this invention, P delays alloying speed | rate. Therefore, the P content is 0.05% or less. P is a solid solution strengthening element, and has the effect of stabilizing ferrite and promoting C concentration in austenite, and the effect of suppressing the generation of red scale in steel to which Si is added. Therefore, the P content is preferably 0.01% or more.

  When S: S is contained in a large amount exceeding 0.01%, the hot ductility is remarkably lowered to induce hot cracking, and the surface properties of the steel sheet are remarkably deteriorated. Moreover, it not only contributes to the strength, but also precipitates as coarse MnS, reducing the ductility such as hole expandability. Therefore, the S content is 0.01% or less. The smaller the amount, the better. However, from the viewpoint of improving the hole expansion property, 0.005% or less is more preferable.

Al: Al is a ferrite stabilizing element, and promotes recrystallization of processed austenite to greatly increase the Ar ₃ transformation point of steel, thereby suppressing the development of crystal orientation necessary for increasing the Young's modulus. Moreover, when it contains abundantly, an austenite single phase area | region will lose | disappear, and it will become difficult to complete | finish rolling in an austenite area | region at the time of hot rolling. Therefore, the Al content is 1.5% or less. In addition, Al has the effect of promoting ferrite formation in the cooling process during annealing, concentrating C in austenite to stabilize austenite, and promoting the formation of low-temperature transformation phase, so the Al amount is 0.2% It is desirable to set it above.

  If N: N is contained in a large amount exceeding 0.01%, slab cracking may be induced during hot rolling, and the surface properties of the steel sheet may be deteriorated. Therefore, the N content is 0.01% or less.

  Nb: Nb is the most important element in the present invention. That is, Nb suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves the Young's modulus after annealing. In addition, the recrystallization of the processed ferrite is suppressed during the temperature rising process during annealing, the austenite transformation from the unrecrystallized ferrite is promoted, and the orientation of the ferrite advantageous for increasing the Young's modulus is developed in the subsequent cooling process. Furthermore, it precipitates as fine carbonitride and contributes to an increase in strength. In order to obtain such an effect, the amount needs to be 0.01% or more, preferably 0.03% or more. On the other hand, if the amount exceeds 0.40%, all of the carbonitride cannot be dissolved at the time of reheating the slab, and coarse carbonitride remains, so recrystallization of processed austenite is suppressed during hot rolling. In addition, since recrystallization of the processed ferrite is not suppressed even during annealing after cold pressing, it is not possible to develop a ferrite orientation that is advantageous for increasing the Young's modulus. Even when the slab is directly hot-rolled after continuous casting, recrystallization of the processed austenite is not suppressed. Therefore, the Nb content is 0.01 to 0.40%, preferably 0.03 to 0.40%. In addition, Nb nitride precipitates at a higher temperature than Nb carbide, so it tends to be coarse during hot rolling, so it has the effect of suppressing recrystallization of processed austenite and suppressing recrystallization of processed ferrite during annealing. Small and has little contribution to increasing Young's modulus. Therefore, since it is necessary to secure the amount of Nb precipitated as carbide, it is necessary to satisfy [Nb] − (93/14) × [N] ≧ 0.005.

  The balance is Fe and inevitable impurities.

  In addition to the above component elements, it is preferable to contain at least one element of V: 0.01 to 0.5% and Ti: 0.01 to 0.2% by mass% for the following reasons.

  V: V precipitates as fine carbonitride and contributes to an increase in strength. For that purpose, the amount needs to be 0.01%. On the other hand, if the amount exceeds 0.5%, the effect of increasing the strength is small and the cost is increased. Therefore, the V amount is 0.01 to 0.5%.

  Ti: Ti precipitates as fine carbonitrides and contributes to an increase in strength. Also, recrystallization of processed austenite during hot rolling is suppressed, and ferrite transformation from unrecrystallized austenite is promoted, thereby contributing to a higher Young's modulus. In order to have such an effect, the amount needs to be 0.01% or more. On the other hand, if the amount exceeds 0.2%, the carbonitride cannot completely dissolve when the slab is reheated, and coarse carbonitride remains, so that the strength is increased or recrystallization of the processed austenite is performed. The effect of suppressing cannot be obtained. Even when the slab is directly hot-rolled after continuous casting, the effect of suppressing recrystallization of the processed austenite is small, resulting in an increase in cost. Therefore, the Ti amount is set to 0.01 to 0.2%.

  Further, for the following reasons, at least one element of Mo: 0.1-1.0%, Cr: 0.1-1.0%, Ni: 0.1-1.0%, B: 0.0005-0.0020%, or Cu: 0.1- It is preferable to contain 2%.

  Mo: Mo is an element that increases the hardenability by reducing the mobility of the interface, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an action, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the effect is not only saturated but also the cost is increased. Therefore, the Mo amount is 0.1 to 1.0%.

  Cr: Cr is an element that suppresses the formation of cementite and enhances hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an effect, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the effect is not only saturated but also the cost is increased. Therefore, the Cr content is 0.1 to 1.0%. When hot dip galvanizing is applied to the steel sheet of the present invention, the Cr oxide produced on the surface induces non-plating, so the Cr content is more preferably 0.5% or less.

Ni: Ni is an element that enhances hardenability by stabilizing austenite, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing and greatly contributes to high strength. In addition, it lowers the Ac ₁ transformation point in the temperature raising process during annealing, promotes austenite transformation from unrecrystallized ferrite, and develops a ferrite orientation that is advantageous for increasing the Young's modulus in the cooling process. In addition, it suppresses recrystallization of processed austenite during hot rolling, lowers the Ar ₃ transformation point, and enables hot rolling in a lower temperature range, thereby promoting ferrite transformation from unrecrystallized austenite. To develop the {113} <110> orientation and improve the Young's modulus after annealing. Furthermore, cracks during hot rolling that occur when Cu is added are prevented. In order to obtain such an action, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the formation of ferrite necessary for increasing the Young's modulus during the cooling process during annealing is suppressed, or the low-temperature transformation phase is increased to increase the strength of the steel and formability. Reduce. Therefore, the Ni content is 0.1 to 1.0%.

  B: B is an element that suppresses the transformation from austenite to ferrite and improves the hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an effect, the amount needs to be 0.0005% or more. On the other hand, if the amount exceeds 0.0020%, the formation of ferrite is remarkably suppressed during the cooling process during annealing, and the Young's modulus is lowered. Therefore, the B amount is 0.0005 to 0.0020%.

  Cu: Cu is an element that enhances hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to high strength. In order to obtain this effect, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 2.0%, the hot ductility is lowered to induce cracking, and the effect of hardenability is saturated. Therefore, the Cu content is 0.1 to 2.0%.

2) Microstructure Column-shaped structural members such as automobile center pillars must have excellent formability such as elongation. Therefore, it is necessary to increase the strength by a composite structure including a ferrite phase and a low-temperature transformation phase such as martensite or bainite. At this time, in order to obtain TS of 590 MPa or more, the volume ratio of the low temperature transformation phase needs to be 1 to 60%. On the other hand, when the volume ratio of the low temperature transformation phase exceeds 60%, the strength is excessively increased and the moldability such as elongation is remarkably lowered.

  Here, the volume ratio of the low-temperature transformation phase is present in a 100 mm square region by observing the position of the plate thickness 1/4 of the plate thickness cross section parallel to the rolling direction of the steel plate with an optical microscope or a scanning electron microscope at a magnification of 1000. The area ratio occupied by the low-temperature transformation phase to be obtained was determined by image processing and used as the volume ratio.

3) X
In general, the rigidity of a certain member is expressed by the product of the Young's modulus E of the material and the moment of inertia of the section, and the moment of inertia of the section is proportional to the λth power of the thickness t of the material. It can be expressed as equation (2).
Rigidity = A × E × t ^λ (2)
In column-shaped structural members such as automobile center pillars, λ is close to 1, and the normal steel sheet E is 200 to 210 GPa. Therefore, even if the sheet thickness is reduced by 10% to reduce weight, the rigidity is constant. In order to maintain this, E must be improved by 11% to 230 GPa.

  E is the slope of the straight line portion in the low strain region of the stress-strain curve obtained by the tensile test shown in Fig. 1, but as the plate thickness is decreased, higher stress is applied to the material. It is not sufficient to evaluate the stiffness with E, and it is necessary to evaluate with the slope of the stress-strain curve near the stress actually applied to the material. In particular, high-strength cold-rolled steel sheets with a low-temperature transformation phase are thought to have many movable dislocations, but the stress-strain curve slope decreases from the low-strain range, ensuring sufficient rigidity as described above. Inability to do so causes problems such as bending and twisting.

  Therefore, the present inventors investigated how much stress is applied from the outside when a high-strength cold-rolled steel sheet having a TS of 590 MPa or more is applied to a columnar structural member. It became clear that it was loaded. Therefore, if 40% stress of TS (approx. 240 MPa) is applied, that is, if the stress-strain curve slope is set to 230 GPa or more when stress around 0.40 × TS is applied, the plate thickness will be 10% or more. Even if it decreases, the rigidity of the columnar structural member can be secured.

  Since it is difficult to determine the slope of the stress-strain curve with high accuracy, as shown in FIG. 1, the slopes of two points of 0.40 × TS and 0.45 × TS were calculated using the above equation (1). It was evaluated with. That is, if X is 230 GPa or more, the rigidity of the columnar structural member can be secured. If X is 230 GPa or more, E is inevitably 230 GPa or more. Further, if the slope near 0.60 × TS is 230 GPa or more, rigidity can be ensured even when higher stress is applied.

  On the surface of the high-strength cold-rolled steel sheet of the present invention, a plated layer of pure zinc, zinc-based alloy, pure Al, Al-based alloy or the like can be provided by electroplating or hot dipping.

4) Manufacturing method The basic idea in the manufacturing method which is an example of the present invention is that in order to set X to 230 GPa or more, first, E is increased to ensure 230 GPa or more, and then stress around 0.4 × TS is applied. It is to improve the slope of the stress-strain curve, that is, X.

4-1) Hot rolling conditions After the slab has been melted in a converter or electric furnace, the total reduction at 950 ° C or lower is 30% or higher, and the hot rolling is performed at a finishing temperature of Ar ₃ transformation point to 900 ° C. When rolled, unrecrystallized austenite consisting of {112} <111> orientation develops immediately after hot rolling, and in the subsequent cooling process, this unrecrystallized austenite transforms into ferrite of {113} <110> orientation, During annealing after cold rolling, ferrite of {112} <110> orientation is generated, and Young's modulus can be improved.

4-2) Winding temperature When winding the steel sheet after hot rolling, if the winding temperature exceeds 650 ° C, the Nb carbonitrides will become coarse and ferrite recrystallization will occur during the heating process during annealing. The effect of suppressing becomes small, and it becomes difficult to transform from non-recrystallized ferrite to austenite. As a result, the orientation of ferrite that transforms in the subsequent cooling process cannot be controlled, and the Young's modulus is greatly reduced. Therefore, the coiling temperature is 650 ° C. or less. Note that if the coiling temperature is too low, the shape is likely to be defective, and therefore, the temperature is preferably 350 ° C. or higher.

4-3) Reduction ratio By cold rolling the hot-rolled steel sheet, the {113} <110> orientation developed in the hot-rolled steel sheet can be rotated to the {112} <110> orientation. During annealing, ferrite with {112} <110> orientation can be generated, and Young's modulus can be increased. In order to obtain such an effect, the rolling reduction during cold rolling needs to be 40% or more.

4-4) Heating conditions during annealing Although the steel sheet after cold rolling is annealed, the ferrite with {112} <110> orientation is transformed to austenite during heating during annealing, and then the austenite is transformed into {112 } In order to reversely transform to ferrite with <110> orientation and improve Young's modulus, it is necessary to transform a sufficient amount of ferrite to austenite during heating. For this purpose, the heating temperature must be (875-400 × [C] -30 × [Mn]) ° C. or higher. On the other hand, if the heating temperature is too high, austenite becomes coarse, and the accumulation of reverse transformed ferrite in the {112} <110> orientation decreases. Therefore, the heating temperature needs to be (915-400 × [C] -30 × [Mn]) ° C. or less. Here, the temperature from (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn]) ° C. represents the temperature range near the Ac ₃ transformation point. . That is, in the steel of the present invention, the Ac ₃ transformation point is represented by (895-400 × [C] -30 × [Mn]) (empirical formula obtained by the inventors), and the above temperature range is approximately the Ac ₃ transformation point. Represents a range of ± 20 ° C. The heating time is 1 s or more in order to promote the transformation from ferrite to austenite. On the other hand, if it is too long, austenite becomes coarse, so the heating time needs to be 600 s or less. Furthermore, in order to suppress the coarsening of austenite and improve the Young's modulus, or to refine the structure and increase the strength, the heating time is preferably set to 300 s or less.

4-5) Cooling conditions during annealing Heated steel sheets must be retained for at least 20 s in the temperature range of 700-500 ° C in order to sufficiently generate ferrite with the {112} <110> orientation during the cooling process. There is. On the other hand, if the residence time is long, carbides are generated and it is difficult to produce a low temperature transformation phase, and the strength tends to decrease. Therefore, the residence time is preferably 200 s or less. The residence time can be controlled by changing the cooling rate in the above temperature range or keeping the temperature constant.

4-6) Tempering conditions during annealing The steel sheet retained in the above temperature range was further tempered by retaining for 60 s or more in the temperature range of 200 to 500 ° C, or cooled to 200 ° C or less after retention in the above temperature range. After that, reheat and stay in the temperature range of 200-500 ° C for 60s or more and temper it to reduce the mobile dislocations that exist in large amounts around the low-temperature transformation phase, or as a solid solution C, N in the mobile dislocations If the interstitial elements are precipitated to prevent dislocation movement, the above-mentioned X can be increased to 230 GPa or more. At this time, if the tempering temperature is less than 200 ° C., the effect cannot be obtained. On the other hand, if the temperature becomes too high and exceeds 500 ° C., the low temperature transformation phase is remarkably softened, so that the temperature range of 200 to 500 ° C. is over 60 s. It must be tempered and retained. Preferably, it is necessary to retain for 60 seconds or more in a temperature range of 200 to 450 ° C. If the residence time is less than 60 seconds, it is difficult to sufficiently obtain the above effect. Further, if the residence time is long, the carbide is coarsened, the low temperature transformation phase is remarkably softened, and the strength is lowered. Therefore, the residence time is preferably 1 hour or less. In addition, residence for 60 seconds or more can be performed by adjusting the heating rate and the cooling rate, or by keeping the temperature constant.

  Steel A having the components shown in Table 1 was melted in a laboratory in a vacuum melting furnace and subjected to hot rolling, cold rolling, and annealing to produce a cold rolled steel sheet. At this time, heating conditions prior to hot rolling: 1 hour at 1250 ° C., total rolling reduction of 950 ° C. or lower in hot rolling: 40%, finishing temperature: 860 ° C., plate thickness after hot rolling: 3.0 mm, winding Rolling conditions: Holding at 600 ° C for 1 hour and then furnace cooling (winding temperature 600 ° C), cold rolling reduction: 50%, plate thickness after cold rolling: 1.5mm, heating during annealing Conditions: 180 s at 820 ° C, average cooling rate from 700 to 500 ° C: 8 ° C / s (residence time: 25 s), average cooling rate from 500 to 300 ° C: 15 ° C / s, tempering conditions: 180 s at 300 ° C Then, after cooling to 200 ° C at an average cooling rate of 10 ° C / s, air cooling was the basic condition, and the cold rolling reduction, annealing heating temperature, 700-500 ° C residence time, and tempering temperature and time were changed. . That is, the above conditions are the same except for the changed conditions. In addition, the temperature range of 700 to 500 ° C. was cooled at an average cooling rate of 8 ° C./s, then air-cooled to room temperature, and reheated to change the tempering temperature and time.

  Then, a 10 × 60 mm test piece was cut out from the produced cold-rolled steel sheet in a direction perpendicular to the rolling direction, and Young's modulus E was measured by a resonance method. In addition, JIS No. 5 tensile specimens were cut from a direction perpendicular to the rolling direction from cold rolled steel sheets subjected to temper rolling at 0.5% after annealing, and tensile properties (TS and elongation El) were measured. Furthermore, from a cold rolled steel sheet not subjected to temper rolling, a JIS No. 5 tensile test piece was cut out from a direction perpendicular to the rolling direction, stress was applied by a tensile test, and 0.40 × TS and 0.45 × using a strain gauge. The amount of strain at the time of stress loading of TS and 0.60 × TS and 0.65 × TS was measured, and X (0.4TS) and X (0.6TS) were obtained from the above formula (1). Furthermore, the volume fraction of the low temperature transformation phase was determined by the method described above. The phases other than the low temperature transformation phase were ferrite phases.

  The cold-rolled steel sheet produced under the basic conditions has TS: 790 MPa, El: 20%, E: 245 GPa, X: 243 GPa, and exhibits an excellent strength-ductility balance and high rigidity.

  Table 2 shows the results obtained by changing the cold rolling reduction ratio. At a rolling reduction of 40% or more according to the present invention, an excellent balance between strength and ductility is exhibited, and X is 230 GPa or more and high rigidity is obtained.

  Table 3 shows the results obtained by changing the heating temperature during annealing. The heating temperature of (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn]) ° C. according to the present invention shows an excellent strength-ductility balance. , X (0.4TS) and X (0.6TS) are both 230 GPa or more, and high rigidity is obtained.

  Table 4 shows the results obtained by changing the residence time at 700 to 500 ° C. The residence time was changed by changing the cooling rate. In the residence time of 20 s or more according to the present invention, an excellent balance between strength and ductility is exhibited, and both X (0.4 TS) and X (0.6 TS) are 230 GPa or more, and high rigidity is obtained.

  Table 5 shows the results obtained by changing the tempering temperature and time when tempering without reheating. In addition, from 500 degreeC to the tempering temperature, it cooled with the average cooling rate of 15 degree-C / s, and after tempering with the tempering temperature and time as described in a table | surface, it cooled to 200 degreeC with the average cooling rate of 10 degree-C / s, and air-cooled. Here, when tempering at 100 ° C. and 200 ° C., air cooling was performed after tempering. When the tempering temperature of the present invention is 200 to 500 ° C. and the residence time in this temperature range is 60 s or more, an excellent balance between strength and ductility is exhibited, and X (0.4TS) and X (0.6 TS) are Both are over 230 GPa and high rigidity is obtained.

  Table 6 shows the results obtained by changing the tempering temperature and time when reheating and tempering. It was heated at an average temperature of 10 ° C./s until the tempering temperature, tempered at the tempering temperature and time shown in the table, then cooled to 200 ° C. at an average cooling rate of 10 ° C./s and air-cooled. Here, when tempering at 100 ° C. and 200 ° C., air cooling was performed after tempering. When the tempering temperature of the present invention is 200 to 500 ° C. and the residence time in this temperature range is 60 s or more, an excellent balance between strength and ductility is exhibited, and X (0.4TS) and X (0.6 TS) are Both are over 230 GPa and high rigidity is obtained.

  Steels B to Y having the components shown in Table 7 were melted in a laboratory in a vacuum melting furnace and subjected to hot rolling, cold rolling, and annealing under the basic conditions of Example 1 to cold rolled steel plates 1 to 24. Was made. At this time, the heating temperature at the time of annealing was changed according to the amounts of C and Mn, and was within the scope of the present invention. Then, the same investigation as in Example 1 was performed. As in Example 1, the phases other than the low temperature transformation phase were ferrite phases.

  The results are shown in Table 8. In the steel plates 1 to 10, 12 to 15, 18, and 21 to 24 that satisfy the component conditions of the present invention, an excellent strength-ductility balance is exhibited, and X (0.4TS) is 230 GPa or more and high rigidity is obtained. On the other hand, [Nb] − (93/14) × [N] is a steel plate 11 outside the scope of the present invention, steel plates 16 and 17 whose Mn amount is outside the scope of the present invention, and steel plates 19 and 20 whose C amount is outside the scope of the present invention. Then, E is low, X (0.4TS) is less than 230GPa, and rigidity is low. Further, with the steel plate 20 having a remarkably low amount of C, TS is 450 MPa and high strength cannot be obtained.

It is a figure explaining the method to obtain | require X.

Claims (6)

In mass%, C: 0.02 to 0.20%, 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, Nb: Made of 0.01 to 0.40%, balance Fe and inevitable impurities, satisfying [Nb]-(93/14) × [N] ≧ 0.005, obtained by tempering by staying in the temperature range of 200-500 ° C for 60s or more It has a microstructure composed of a ferrite phase and a low temperature transformation phase, the volume ratio of the low temperature transformation phase is 1 to 60%, the tensile strength is 590 MPa or more, and X in the following formula (1) is 230 GPa or more. High strength cold-rolled steel sheet with high rigidity;
X = 0.005 × TS / (ε2-ε1) [GPa] (1)
Here, [M] represents the content of element M [% by mass], TS is the tensile strength [MPa] of the steel sheet, ε1 is the nominal strain [%] when the nominal stress is 0.40 × TS, and ε2 is the nominal stress Represents the nominal strain [%] when is 0.45 × TS.   2. The high-strength cold-rolled steel sheet having high rigidity according to claim 1, wherein the mass percentage is Nb: 0.03 to 0.40%.   The high-strength cold-rolled steel sheet having high rigidity according to claim 1 or 2, further comprising at least one element of V: 0.01 to 0.5% and Ti: 0.01 to 0.2% by mass%.   Further, according to any one of claims 1 to 3, further comprising at least one element of Mo: 0.1 to 1.0%, Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, B: 0.0005 to 0.0020% in mass%. A high-strength cold-rolled steel sheet having high rigidity according to item 1.   The high-strength high-strength cold-rolled steel sheet according to any one of claims 1 to 4, further comprising Cu: 0.1 to 2.0% by mass. A steel slab having the composition according to claims 1 to 5, wherein the total reduction amount at 950 ° C or lower is 30% or more, hot-rolled at a finishing temperature of Ar ₃ transformation point to 900 ° C, and wound at 650 ° C or lower After cold rolling at a rolling reduction of 40% or more, heat for 1 to 600 s at (875-400 x [C] -30 x [Mn]) to (915-400 x [C] -30 x [Mn]) ° C And then cooled and allowed to stay for 20 s or more in the temperature range of 700 to 500 ° C. and tempered by staying in the temperature range of 200 to 500 ° C. for 60 s or after staying in the temperature range of 700 to 500 ° C. Cooling to 200 ° C or lower, reheating, staying in the temperature range of 200-500 ° C for 60s or longer and annealing under conditions of tempering; manufacturing method of high strength cold-rolled steel sheet with high rigidity; where [M] is an element M content [% by mass] is expressed.

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