CN1386140A - Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same - Google Patents

Abstract

An object of the present invention is to provide a cold-rolled steel sheet and an alloyed hot-dip galvanized steel sheet in which tensile strength is effectively increased by press forming and heat treatment while maintaining excellent deep drawability in press forming. Specifically, a steel composition contains less than 0.01% of C, 0.005 to 1.0% of Si, 0.01 to 1.0% of Mn, 0.005 to 0.050% of Nb, 0.005 to 0.030% of Al, 0.005 to 0.040% of N, 0.0005 to 0.0015% of B, 0.05% or less of P, and 0.01% or less of S, the balance substantially composed of Fe, in which the following equations (1) and (2) are satisfied: N% >=0.0015 + 14/93-Nb% + 14/27-Al% + 14/11-B% C% <= (12/93)-Nb%.

Description

Cold rolled steel sheet having strain age hardening property, galvanized steel sheet and manufacturing method thereof

technical field

The present invention relates to parts that require structural strength, especially strength and/or rigidity during deformation, such as building components, mechanical structural parts, and automotive structural parts. A cold-rolled steel sheet, an electrogalvanized steel sheet, a hot-dip galvanized steel sheet, and an alloyed hot-dip galvanized steel sheet, which are suitable as a raw material steel sheet for a formed body and have excellent strain-age hardening properties, and methods for producing the same.

In addition, in the present invention, the term "excellent in strain age hardening properties" means that after pre-deformation with a tensile strain of 5%, aging treatment is performed at a temperature of 170° C. for 20 minutes, before and after the aging treatment. The increase in deformation stress (denoted as BH amount; BH amount=yield stress after aging treatment-pre-deformation stress before aging treatment) is more than 80MPa, and the stretching before and after strain aging treatment (predeformation + aging treatment) The strength increase (denoted as ΔTS; ΔTS=tensile strength after aging treatment-tensile strength before pre-deformation) is above 40 MPa.

Background technique

When manufacturing a press-formed steel sheet, there is a method of spraying and baking at less than 200°C as a method of forming a soft material before stamping to facilitate stamping, and hardening it after stamping to improve the strength of the part. The steel plate used for painting and baking has developed BH steel plate.

For example, Japanese Patent Application Laid-Open No. 55-141526 discloses adding Nb according to the C, N, and Al contents in the steel, while limiting Nb/(solid solution C+solid solution N) represented by at% to a specific range, by controlling The cooling rate after annealing is used to adjust the solid solution C and solid solution N in the steel sheet; Japanese Patent Publication No. 61-45689 discloses a method of improving bake hardenability by adding Ti and Nb in combination.

However, in order to form the above-mentioned steel sheet with excellent deep drawability, the strength of the raw material steel sheet is low, and it is not necessarily sufficient as a structural material.

In addition, JP-A-5-25549 discloses a method of improving bake hardenability by adding W, Cr, and Mo singly or in combination to steel.

In the prior art described above, the increase in strength by bake hardening is based on the effect of a small amount of solid solution C and solid solution N in the steel plate. In addition, as is well known, in the case of BH steel plate, only the yield of the material is increased. An increase in strength does not increase tensile strength. Therefore, only the effect of increasing the deformation initial stress of the part, and the effect of increasing the stress (tensile strength after forming) required for deformation of the entire deformation region from the beginning of deformation to the completion of deformation cannot be said to be sufficient.

As a cold-rolled steel sheet with increased tensile strength after forming, for example, JP-A-10-310847 discloses an alloyed hot-dip galvanized steel sheet with an increased tensile strength of 60 MPa or more in a heat treatment temperature range of 200 to 450°C.

The composition of this steel sheet is represented by mass%, and contains C: 0.01 to 0.08%, Mn: 0.01 to 3.0%, and contains one or more of W, Cr, and Mo in a total amount of 0.05 to 3.0%. One or two or more of Ti: 0.005-0.1%, Nb: 0.005-0.1%, and V: 0.005-0.1% are contained as needed, and the microstructure of steel is composed of ferrite or ferrite main body.

However, in this technology, fine carbides are formed in the steel sheet through heat treatment after forming, and the strain imparted during stamping effectively propagates dislocations and increases the amount of strain. Therefore, heat treatment at a temperature range of 220 to 370° C. The biggest difficulty is that the required heat treatment temperature is higher than the general bake hardening treatment temperature.

Furthermore, in connection with the recent regulations on exhaust gas derived from global environmental problems, reduction in vehicle body weight has become an extremely important issue. In order to reduce the body weight of an automobile, it is effective to increase the strength of the steel sheet used, that is, to use a high-strength steel sheet, and to make the steel sheet thinner.

Automobile parts using thin high-strength steel sheets must fully exhibit the characteristics appropriate to their functions. Properties vary depending on the part, for example, static strength corresponding to indentation resistance, bending, torsional deformation, fatigue resistance, impact resistance, etc. That is, high-strength steel sheets used for automobile parts need to be excellent in these characteristics after forming. These characteristics are related to the strength of the steel sheet after forming, so in order to reduce the thickness, it is necessary to set the lower limit of the strength of the high-strength steel sheet to be used.

On the other hand, in the process of manufacturing automotive parts, steel sheets are press-formed. In the case of press forming, when the strength of the steel plate is too high, the problems that arise are: ① Deterioration of shape freezing; ② Decreased ductility, cracks and necking during forming; Indentation resistance (resistance to dents caused by local compressive loads) deteriorates, thereby hindering the spread of application of high-strength steel sheets to automobile bodies.

As a method to break this situation, for example, cold-rolled steel sheets for outer panels are known, which use ultra-low carbon steel as a raw material and control the amount of C remaining in a solid solution state to an appropriate range. This type of steel sheet remains soft during press forming to ensure shape freezeability and ductility, and the yield stress is increased by utilizing the strain age hardening phenomenon caused by the painting and baking process at 170°C for about 20 minutes after press forming to ensure Indentation resistance. This type of steel sheet is soft when C is dissolved in the steel during stamping. On the other hand, after stamping, solid solution C is fixed in the dislocations introduced during stamping during the painting and baking process after stamping, thereby yielding. Stress rises.

However, in such a steel sheet, from the viewpoint of preventing the occurrence of tensile strain that becomes a surface defect, the amount of increase in yield stress due to strain age hardening is kept low. Therefore, there is a difficulty in that the effect of actually contributing to the weight reduction of components is small.

On the other hand, for applications where the appearance is not so much a problem, it has been proposed to use solid solution N to further increase the amount of bake hardening, and to further improve the bake quality by making the structure a composite structure composed of ferrite and martensite. Bake-hardened steel plate.

For example, JP-A-60-52528 discloses that steel containing C: 0.02-0.15%, Mn: 0.8-3.5%, P: 0.02-0.15%, Al: less than 0.10%, and N: 0.005-0.025% is contained in A method of manufacturing a high-strength thin steel sheet that is coiled at a temperature of 550°C or lower and annealed after hot rolling and cold rolling for controlled cooling heat treatment and has good ductility and spot weldability. The steel plate manufactured by the technology described in JP-A No. 60-52528 has a mixed structure composed of ferrite and martensite as the main low-temperature transformation product phase. The strain aging of the N produced paint bake gives high strength.

However, with the technology described in JP-A-60-52528, the increase in yield stress YS caused by strain age hardening is large, but the increase in tensile strength TS is small, and the increase in yield stress YS is also greatly increased. Due to large fluctuations in mechanical properties such as dispersion, it is not possible to expect a thinner steel sheet that contributes to the weight reduction of automobile parts that is currently required.

In addition, Japanese Patent Publication No. 5-24979 discloses that it has a composition containing C: 0.08 to 0.20%, Mn: 1.5 to 3.5%, and the rest is composed of Fe and unavoidable impurities. The structure is composed of 5% ferrite. A bake-hardenable high-strength cold-rolled steel sheet composed of uniform bainite or partially martensitic bainite. In the cold-rolled steel sheet described in Japanese Patent Publication No. 5-24979, in the cooling process after continuous annealing, the temperature range of 400 to 200 ° C is rapidly cooled, and the structure is mainly bainite through the subsequent slow cooling. There is no high amount of bake hardening.

However, although the yield strength of the steel plate described in Japanese Patent Publication No. 5-24979 is increased after painting and baking, and a high bake hardening amount that has not been obtained before is obtained, even the tensile strength cannot be increased, and it cannot be used for strength members. Improvement in fatigue resistance and impact resistance after molding is expected. Therefore, there remains a problem that it cannot be applied to applications that strongly require fatigue resistance properties, impact resistance properties, and the like.

In addition, Japanese Patent Publication No. 61-12008 discloses a method for manufacturing a high-strength steel sheet with a high r value. Annealing in the coexistence zone of celestial body, the obtained steel plate has high r value and high paint bake hardenability (BH property), but the obtained BH amount is at most about 60MPa. In addition, although the yield point of the steel plate also increases after aging, However, TS did not increase, and there was a problem that applicable parts were limited.

Furthermore, although the above-mentioned conventional steel sheets are excellent in strength evaluation after painting and baking by a simple tensile test, the strength at the time of plastic deformation varies greatly depending on actual stamping conditions. It may not be sufficient for parts requiring reliability.

Among the painted and baked steel sheets of press-formed products, regarding hot-rolled steel sheets, for example, Japanese Patent Publication No. 8-23048 discloses that the tensile strength is soft at the time of processing, and is effective for improving fatigue properties through the baking treatment after processing. Substantial rise in the manufacturing method of hot-rolled steel sheets.

In this technique, the amount of C is set to 0.02 to 0.13% by mass, and a large amount of N is added to 0.0080 to 0.0250% by mass. In addition, the finish rolling temperature and coiling temperature are controlled so that a large amount of solid solution N remains in the steel, and the metal The structure is a composite structure mainly composed of ferrite and martensite, so that the increase in tensile strength of more than 100 MPa can be achieved at the post-forming heat treatment temperature: 170°C.

In addition, Japanese Patent Application Laid-Open No. 10-183301 discloses that among the steel components, especially, C and N are limited to C: 0.01 to 0.12% by mass and N: 0.0001 to 0.01% by mass, while controlling the average grain size to 8 μm or less. According to this, it is possible to ensure a high BH content of 80 MPa or more, and at the same time suppress the AI content to 45 MPa or less, which is a hot-rolled steel sheet with excellent bake hardenability and room temperature aging resistance.

However, since these steel sheets are hot-rolled sheets, the austenite/ferrite transformation after finish rolling makes the ferrite structure random, so it is difficult to obtain a high r value, and it is difficult to say that it has sufficient deep-drawing strength. Ductility.

Moreover, even if the hot-rolled steel sheet obtained by these techniques is used as the raw material, even if it is cold-rolled and recrystallized annealed, it is not necessarily possible to obtain the same increase in tensile strength after forming-heat treatment and a high BH of 80 MPa or more as the hot-rolled steel sheet. This is because: the steel microstructure becomes different from hot rolling through cold rolling and recrystallization annealing. In addition, large strain accumulation is caused during cold rolling, so it is easy to form carbides, nitrides or carbonitrides, solid solution C and The state of solid solution N changes.

The present invention has been developed based on the above-mentioned actual situation, and its object is to provide a cold-rolled steel sheet and a steel sheet excellent in strain age hardening characteristics in which tensile strength is increased by press forming-heat treatment while maintaining excellent deep drawability during press forming and heat treatment. Alloyed hot-dip galvanized steel sheets and their advantageous manufacturing methods.

In addition, in view of the above-mentioned problems of the prior art, an object of the present invention is to provide a deep-drawn steel sheet with a TS×r value of ≥750 MPa and an excellent deep-drawing property and an excellent strain-age hardening property (BH≥80 MPa and ΔTS≥40 MPa). Cold-rolled steel sheets and hot-dip galvanized steel sheets (also including alloyed hot-dip galvanized steel sheets) and their advantageous manufacturing methods.

Further, the object of the present invention is also to solve the above-mentioned problems of the prior art, and to provide an automobile component suitable for automobile parts requiring a high degree of formability, which has high formability and stable quality characteristics when it is soft, and is easy to form into complex parts. Shaped auto parts, there is no springback, torsion, warpage and other shape defects, cracks, etc., and the heat treatment of auto parts after being formed into auto parts can obtain sufficient strength, which can fully contribute to the development of the car body. Lightweight, high-strength cold-rolled steel sheets having a high r-value of 1.2 or more and excellent strain-age hardening properties, and a manufacturing method for manufacturing these steel sheets at low cost and without destroying the shape industrially.

Contents of the invention

In order to achieve the above-mentioned problems, the present inventors changed the composition and production conditions variously, produced steel plates, and conducted many material evaluation experiments. As a result, knowledge was obtained that using N, which has not been actively used in the field requiring high workability, as a strengthening element, and making full use of the large strain age hardening phenomenon that appears through the action of this strengthening element, can be easily obtained. Improvement of formability and high strength after forming are achieved at the same time.

Furthermore, the present inventors have found that in order to make full use of the strain age hardening phenomenon caused by N, it is necessary to combine the strain age hardening phenomenon caused by N with the paint baking conditions of automobiles or the heat treatment conditions after positive forming. Therefore, it is effective to make the hot rolling conditions and cold rolling and cold rolling annealing conditions appropriate, and to control the microstructure and solid solution N content of the steel sheet within a certain range. It has also been found that in order to stably develop the strain age hardening phenomenon caused by N, it is important to control the Al content in terms of composition, particularly according to the N content.

Furthermore, the present inventors found that: in order to obtain a high r value, reduce the C content, perform continuous annealing at the temperature of the ferrite-austenite two-phase region, control the subsequent cooling, and form A structure containing more than 5% of the acicular ferrite phase expressed by the area ratio, with the combination of such a microstructure and an appropriate amount of solid solution N, can have a high r value, excellent stamping formability, and strain aging Cold-rolled steel sheet with excellent hardening properties. In addition, it was found that N can be fully and effectively used without causing the problem of aging degradation at room temperature which has been a problem in the past.

That is, the present inventors found that: use N as a strengthening element, control the Al content in an appropriate range according to the N content, and make the hot rolling conditions, cold rolling, and cold rolling and annealing conditions appropriate, so that the microstructure and solid solution N can be optimized. According to this, compared with the conventional solid-solution-strengthened C-Mn steel plate and precipitation-strengthened steel plate, it has a high r value, has particularly excellent formability, and strain age hardening that the above-mentioned conventional steel plate does not have. characteristic steel plate.

In addition, the strength of the steel sheet of the present invention after painting and baking obtained by a simple tensile test is higher than that of conventional steel sheets, and the dispersion of the strength when it is plastically deformed according to actual stamping conditions is small, and stable component strength can be obtained. characteristics, applicable to components requiring reliability. For example, the part where a large strain is applied and the plate thickness is reduced has a larger amount of hardening than other parts. If the load-carrying capacity of (plate thickness) × (strength) is evaluated, it is in the direction of uniformity. The strength of the component is stable.

In order to achieve the above-mentioned purpose, the present inventor further carried out repeated intensive research, and obtained the following knowledge as a result:

1) In order to increase the tensile strength after forming-heat treatment, it is necessary to perform tensile deformation to introduce new dislocations. It is necessary that dislocations introduced by pre-deformation do not migrate even when the upper yield stress is reached by the interaction of dislocations introduced by forming and interstitial elements or precipitates.

2) In order to obtain the above-mentioned interaction by forming carbides, nitrides or carbonitrides of W, Cr, Mo, Ti, Nb, Al, etc., it is necessary to increase the heat treatment temperature after forming to 200°C or higher. Therefore, active effective use of interstitial elements or effective use of Fe carbides or Fe nitrides is advantageous in terms of lowering the heat treatment temperature after forming.

3) Among the interstitial elements, compared with solid solution C, even if the heat treatment temperature after forming is lowered, the interaction between solid solution N and dislocations introduced by forming is greater. Even if the upper yield stress is reached, the dislocations introduced by pre-deformation Dislocations are also difficult to migrate.

4) There are intra-grain and grain boundaries as places where solid-solution N exists in the steel, but the area of the grain boundary is larger in terms of the strength increase after heat treatment after forming. That is, it is advantageous that the crystal grain size is small.

5) From the point of view of enlarging the grain boundary area, while adding Nb and B in combination, the normal growth of ferrite grains after hot rolling can be suppressed by cooling immediately after hot rolling, and it can be suppressed in the following Grain growth in recrystallization annealing after cold rolling.

The present invention has been accomplished based on the above knowledge. The above knowledge is obtained from the following experiments.

Experiment 1

In terms of mass%, it will contain C: 0.0015%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0.008% and N: 0.011%, and in the range of 0.005-0.05%. A thin slab (thickness: 30 mm) containing Nb, Al in the range of 0.005 to 0.03%, and the rest of Fe and unavoidable impurities is uniformly heated at ₁₁₅₀ °C. Hot rolling is carried out at 900°C with a 3-pass type, and after the rolling is completed, it is water-cooled after 0.1 second. Thereafter, a heat treatment equivalent to winding a coil was performed at 500° C. for 1 hour.

The obtained hot-rolled sheet with a thickness of 4 mm was cold-rolled at a reduction rate of 82.5%, then subjected to recrystallization annealing at 800° C. for 40 seconds, and then subjected to temper rolling at a reduction rate of 0.8%. A JIS No. 5 tensile test piece was prepared along the rolling direction from the cold-rolled sheet thus obtained, and the tensile strength was measured at a strain rate of 0.02/sec using a normal tensile testing machine. Also, a JIS No. 5 tensile test piece obtained from these cold-rolled sheets along the rolling direction was given a tensile strain of 10%, heat-treated at 120° C. for 20 minutes, and then used for a normal tensile test. The difference between the tensile strength of the test pieces prepared from these cold-rolled sheets and the test piece heat-treated at 120°C for 20 minutes after applying a tensile strain of 10% was defined as the increase in tensile strength after forming. Quantity (ΔTS).

Fig. 1 shows the results of studies on the relationship between steel components (N%-14/93·Nb%-14/27·Al%-14/11·B%) and ΔTS.

It has been found that, as shown in the figure, when the value of (N%-14/93 Nb%-14/27 Al%-14/11 B%) satisfies 0.0015% by mass or more, ΔTS reaches 60 MPa or more .

Experiment 2

In terms of mass%, it will contain C: 0.0010%, Si: 0.02%, Mn: 0.6%, P: 0.01%, S: 0.009%, N: 0.012%, Al: 0.01% and Nb: 0.015%, and at 0.00005 A thin slab (thickness: 30mm) consisting of B in the range of 0.0025% and Fe and unavoidable impurities in the rest is uniformly heated at 1100°C, and the processing temperature is 920°C above the Ar ₃ transformation point Three-gauge rolling was performed, and after the completion of the rolling, water cooling was performed 0.1 second later, and a heat treatment equivalent to winding a coil was performed at 450° C. for 1 hour.

The obtained hot-rolled sheet with a thickness of 4 mm was cold-rolled at a reduction rate of 82.5%, and then subjected to recrystallization annealing at 820° C. for 40 seconds, followed by temper rolling at a reduction rate of 0.8%.

A JIS No. 5 tensile test piece was prepared along the rolling direction from the cold plate obtained in this way, and the tensile strength was measured at a strain rate of 0.02/sec using a normal tensile testing machine. Also, tensile strains of 10% were applied to tensile test pieces separately obtained from these cold-rolled sheets, and heat-treated at 120° C. for 20 minutes, and then used for ordinary tensile tests.

Fig. 2 shows the results of research on the relationship between the B content in steel and ΔTS. As shown in the figure, it can be seen that when B is contained in an amount of 0.0005 to 0.0015% by mass, a high ΔTS of 60 MPa or more can be obtained.

In addition, by adding Nb and B in combination, the crystal grains are refined, and a high ΔTS can be obtained, which was confirmed by observation of the microstructure.

That is, it is presumed that when the amount of B is less than 0.0005% by mass, the effect of grain refinement by combined addition with Nb is small. On the contrary, when the amount of B exceeds 0.0015% by mass, the amount of B segregated at the grain boundary and its vicinity increases, and since the interaction between B atoms and N atoms is strong, the effective solid-solution N amount decreases, and thus ΔTS decreases.

Experiment 3

In terms of mass%, the composition of steel A is: containing C: 0.0010%, N: 0.012%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0.008%, Nb: 0.014 % and Al: 0.01%, the rest is Fe and unavoidable impurities; the composition of steel B is: containing C: 0.010%, N: 0.0012%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P : 0.03%, S: 0.008%, Nb: 0.014%, and Al: 0.01%, and the remainder is Fe and unavoidable impurities. Each thin slab (thickness: 30mm) of steel A and steel B is uniformly heated at 1150°C, and then 3-pass rolling is performed at a processing temperature of 910°C above the Ar ₃ transformation point. After rolling, After 0.1 second, gas cooling was started, and then heat treatment equivalent to winding a coil was performed at 600° C. for 1 hour.

The obtained hot-rolled sheet with a thickness of 4 mm was cold-rolled at a reduction ratio of 82.5%, and then subjected to recrystallization annealing at 880° C. for 40 seconds, followed by temper rolling at a reduction ratio of 0.8%.

A JIS No. 5 tensile test piece was prepared along the rolling direction from the cold-rolled sheet thus obtained, and the tensile strength was measured at a strain rate of 0.02/sec using a common tensile testing machine. Also, tensile strains of 10% were applied to tensile test pieces separately obtained from these cold-rolled sheets, and heat treatments were performed at various temperatures for 20 minutes, and then used for ordinary tensile tests.

Fig. 3 shows the results of research on the influence of ΔTS on the heat treatment temperature after forming. As shown in the figure, in the relatively low region where the heat treatment temperature after forming is 200°C or lower, Steel A, which is an ultra-low carbon, high-N steel, exhibits a higher temperature than that of a semi-ultra-low carbon and low N steel. Steel B also has high ΔTS, and in the high temperature region, shows the same degree of ΔTS. From these experimental results, it can be seen that it is effective to make full use of solid solution N to ensure ΔTS in the low temperature region.

In addition, Fig. 4 shows the decrease in elongation due to aging at room temperature with respect to grain size d and steel composition (N%-14/93 Nb%-14/27 Al%-14/11 B%) The results of a study on the influence of the amount (ΔE1) and the amount of increase in tensile strength (ΔTS) after forming. Furthermore, the reduction in elongation (ΔE1) was performed as an acceleration of aging at room temperature by using the total elongation measured using a JIS No. 5 test piece obtained from a cold-rolled sheet along the rolling direction and using a test piece obtained separately. The difference in total elongation measured after holding treatment at 100° C. for 8 hours was evaluated.

It can be seen that, as shown in the figure, the value of (N%-14/93.Nb%-14/27.Al%-14/11.B%) is 0.0015% by mass or more, and the grain size d is When the thickness is below 20μm, high ΔTS and low ΔE1 can be obtained at the same time.

Experiment 4

Heat the thin slab of 0.0015%C-0.30%Si-0.8%Mn-0.03%P-0.005%S-0.012%N-0.02～0.08%Al steel uniformly to 1050°C, and then under the premise that the processing temperature is 670°C Hot finish rolling is carried out with 7 passes, followed by recrystallization annealing at 700°C for 5 hours, and the obtained hot-rolled sheet with a thickness of 4mm is cold-rolled at a reduction rate of 82.5%, and then rolled at 875°C for 40 seconds Recrystallization annealing was carried out, followed by temper rolling at a reduction ratio of 0.8%, and JIS No. 5 tensile test pieces were prepared from the obtained cold-rolled sheet, and a normal tensile testing machine was used at a strain rate of 3×10 ^-3 Tensile test is carried out per second, and the TS×r value and ΔTS are measured. The results are shown in Table 5. When N/Al≥0.30 is satisfied, TS×r value≥750MPa and ΔTS≥40MPa can be realized. Furthermore, when N/Al≥0.30, the fact that BH≥80MPa can be achieved was separately confirmed.

Experiment 5

Uniformly heat the thin slab of 0.0015%C-0.0010%B-0.01%Si-0.5%Mn-0.03%P-0.008%S-0.011%N-0.005～0.05%Nb-0.005～0.03%Al steel to 1000℃ , followed by hot finish rolling with 7 passes at a processing temperature of 650°C, and then recrystallization annealing at 800°C for 60 seconds. Cold rolling, followed by recrystallization annealing at 880°C x 40 seconds, then temper rolling at a reduction rate of 0.8%, and JIS No. 5 tensile test pieces were prepared from the resulting cold-rolled sheet, and the usual tensile test pieces were used. Tensile testing was performed with a tensile testing machine at a strain rate of 3×10 ^-3 /sec, and TS×r values, BH, and ΔTS were measured. The relationship between these measured values and N/(Al+Nb+B) is shown in FIG. 5 . In this experiment, using steel containing Nb: 0.005 to 0.05%, B: 0.0010%, as shown in Fig. 5, BH ≥ 80MPa, ΔTS can be realized in the range of N/(Al+Nb+B) ≥ 0.30 ≥60MPa, TS×r value≥850MPa.

Experiment 6

Uniformly heat the thin slab of 0.0010%C-0.02%Si-0.6%Mn-0.01%P-0.009%S-0.015%N-0.01%Al-0.015%Nb-0.0001～0.0025%B steel to 1050°C, then On the premise that the processing temperature is 680°C, the hot rolling is carried out with 7 passes, and then the recrystallization annealing is carried out by intermittent annealing at 750°C for 5 hours. Cold rolling, followed by recrystallization annealing at 880°C × 40 seconds, and then temper rolling at a reduction rate of 0.8%, and JIS No. 5 tensile test pieces were prepared from the cold-rolled plate, and the usual The tensile tester performed a tensile test at a strain rate of 3×10 ^-3 /sec, and measured TS×r values, BH, and ΔTS. The relationship between these measured values and the amount of B is shown in FIG. 6 .

As shown in FIG. 6 , in the range of B: 0.0003 to 0.0015%, in addition to BH ≥ 80 MPa, ΔTS ≥ 60 MPa and TS×r value ≥ 850 MPa, which are higher ΔTS levels than B < 0.0003%, can also be realized. In addition, it can be seen from the observation of the microstructure that the crystal grains are particularly refined within the range of the B amount.

From the results of experiments 5 and 6, it is clear that the range of N/(Al+Nb+B)≥0.30 is stipulated, B≥0.0003%, and Nb is also added in combination, so that the grains are refined, and the value levels of ΔTS and TS×r are further reduced improve. When B<0.0003%, there is no effect of grain refinement by compound addition with Nb. On the other hand, when B>0.0015%, the characteristics are lowered on the contrary. This is presumably because the amount of B segregated at and near the grain boundary increases, and the effective solid-solution N amount decreases due to the strong interaction between B atoms and N atoms. In addition, the same study was conducted on the case of adding Ti and V instead of Nb, and it was confirmed that the same effect as that of Nb can be obtained. The present invention is accomplished based on above knowledge, and its gist is as follows:

The first present invention is a cold-rolled steel sheet excellent in strain age hardening characteristics, characterized by containing C: 0.15% or less, Si: 1.0% or less, Mn: 2.0% or less, and P: 0.1% in mass%. Below, S: 0.01% or less, Al: 0.005 to 0.030%, N: 0.0050 to 0.0400%, and N/Al: 0.30 or more, solid solution N is 0.0010% or more, and the rest is composed of Fe and unavoidable impurities composition.

In the first invention, among the aforementioned compositions, the following ranges are particularly preferable. That is, a cold-rolled steel sheet excellent in strain age hardening characteristics, characterized by containing C: less than 0.01%, Si: 0.005-1.0%, Mn: 0.01-1.5%, and P: 0.1% or less in mass %. , S: 0.01% or less, Al: 0.005 to 0.030%, N: 0.005 to 0.040%, and N/Al: 0.30 or more, solid solution N is 0.0010% or more, and the rest is composed of Fe and unavoidable impurities composition.

In the first invention, in addition to the above composition, B: 0.0001 to 0.0030% and Nb: 0.005 to 0.050% in mass % are preferably contained within the ranges satisfying the following formulas (1) and (2).

N%≥0.0015+14/93 Nb%+14/27 Al%+14/11 B% ---(1)

C%≤0.5·(12/93)·Nb% ---(2)

In the first invention, in addition to the above composition, if necessary, one or two or more of Cu, Ni, and Mo may be contained in a total amount of 1.0% or less in mass %.

In the first invention, the crystal grain size of the steel sheet is preferably 20 μm or less.

In the first invention, in the heat treatment temperature: 120 to 200° C. low-temperature range, there is an increase in strength after forming: 60 MPa or more.

In the first invention, electrogalvanized, hot-dip galvanized, and alloyed hot-dip galvanized layers may be provided on the surface of the above-mentioned cold-rolled steel sheet.

The second present invention is a method for producing a cold-rolled steel sheet excellent in strain age hardening characteristics, characterized in that: expressed in mass%, C: less than 0.01%, Si: 0.005-1.0%, and Mn: 0.01-1.5% , P: 0.1% or less, S: 0.01% or less, Al: 0.005-0.030%, N: 0.005-0.040%, and the content range satisfies N/Al: 0.30 or more, and the rest is substantially Fe. Rolling, at this time, start cooling immediately after finishing rolling, coil at a coiling temperature: 400-800°C, and then perform cold rolling at a reduction ratio: 60-95%, and then perform at a temperature of 650-900°C Recrystallization annealing.

The second invention preferably contains B: 0.0001-0.0030% and Nb: 0.005-0.050% expressed in mass % in addition to the composition described above, within the ranges satisfying the following formulas (1) and (2).

N%≥0.0015+14/93 Nb%+14/27 Al%+14/11 B% ---(1)

C%≤0.5·(12/93)·Nb% ---(2)

In the second invention, it is preferable to raise the temperature at a rate of 1 to 20° C./sec in the temperature range from 500° C. to the recrystallization temperature during the temperature raising process of the above-mentioned recrystallization annealing.

In the second invention, after recrystallization annealing, hot-dip galvanizing treatment may be performed, followed by heat alloying treatment.

The third present invention is a cold-rolled steel sheet for deep drawing with excellent strain age hardening characteristics, characterized by containing C: 0.01% or less, Si: 1.0% or less, and Mn: 0.01 to 1.5% in mass%. , P: 0.1% or less, S: 0.01% or less, Al: 0.005-0.020%, N: 0.0050-0.040%, and N/Al: 0.30 or more, solid solution N is 0.0010% or more, and the rest is composed of Fe and Composition composed of unavoidable impurities, its TS×r value: 750MPa or more.

In the third invention, in addition to the above composition, B: 0.0001 to 0.0030% and Nb: 0.005 to 0.050% are preferably contained in ranges satisfying the following formulas (1) and (2).

N%≥0.0015+14/93 Nb%+14/27 Al%+14/11 B% ---(1)

C%≤0.5·(12/93)·Nb% ---(2)

In the third invention, in addition to the above composition, B: 0.0001 to 0.0030%, Nb: 0.005 to 0.050%, Ti: 0.005 to 0.070%, and V: 0.005 to 0.10% are contained in mass %. 1 or 2 or more components, preferably N/(Al+Nb+Ti+V+B): 0.30 or more, and solid solution N: 0.0010% or more.

The fourth present invention is a method of manufacturing cold-rolled steel sheets for deep drawing with excellent strain age hardening characteristics, characterized in that: expressed in mass %, C: 0.01% or less, Si: 0.005 to 1.0%, Mn : 0.01～1.0%, P: below 0.1%, S: below 0.01%, Al: 0.005～0.030%, N: 0.005～0.040%, including B: 0.0003～0.0030%, Nb: 0.005～0.050%, Ti: 0.005 ～0.070%, V: 0.005～0.10%, one or two or more components, and N/(Al+Nb+Ti+V+B): 0.30 or more. Rough rolling is carried out at a finishing temperature of 1000°C or lower and Ar ₃ or higher, followed by finish rolling and coiling while lubricating in the temperature range of Ar ₃ or lower and 600°C or higher. The total reduction ratio is more than 80%, and the obtained hot-rolled sheet is subjected to recrystallization annealing, followed by cold rolling at a reduction ratio of 60-95%, and the obtained cold-rolled sheet is subjected to recrystallization annealing.

The fifth present invention is a high-strength cold-rolled steel sheet excellent in formability, strain age hardening properties, and aging resistance at room temperature, characterized in that it contains C: 0.0015% to 0.025%, and Si: 1.0% or less, expressed in mass%. , Mn: 2.0% or less, P: 0.1% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.005 0-0.0250%, and B: 0.0005-0.0050%, Nb: 0.002-0.050% of 1 One or two or more kinds of components, and make N/Al 0.3 or more, solid solution N 0.0010% or more, and the rest is composed of Fe and unavoidable impurities, and the needle is composed of 5% or more in area ratio. Structure composed of ferrite phase and average particle size: 20 μm or less, and its r value: 1.2 or more.

In the fifth aspect of the present invention, in addition to the aforementioned composition, it is preferable to further contain one or more components of the following groups a to c represented by mass %, wherein,

Group a: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group b: one or two of Ti and V, the total amount is less than 0.1%;

Group c: 1 or 2 kinds of Ca and REM, the total amount is 0.0010-0.010%.

The sixth present invention is a method for manufacturing a high-strength cold-rolled steel sheet having an r value of 1.2 or more, excellent formability, strain age hardening properties, and aging resistance at room temperature, characterized in that: expressed in mass %, C: 0.0015-0.025%, Si: 1.0% or less, Mn: 2.0% or less, P: 0.1% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050-0.0250%, and B: 0.0003-0.0050% , Nb: 0.002 to 0.050% of one or more kinds of steel slabs, and the composition of N/Al of 0.3 or more is heated to the slab heating temperature: 1000°C or more, rough rolling is performed to form a thin slab, and the sheet is sequentially subjected to The billet is subjected to finish rolling at a temperature of more than 800°C, and coiled at a coiling temperature of less than 650°C to form a hot-rolled sheet. The hot-rolled sheet is pickled and cold-rolled to form a cold-rolled sheet. The cold-rolling process of the rolled plate, the cold-rolled plate is continuously annealed at the temperature in the ferrite-austenite two-phase region, and cooled to the temperature zone below 500°C at a cooling rate of 10-300°C/sec. Cold-rolled sheet annealing process.

In the sixth aspect of the present invention, in addition to the aforementioned composition, it is preferable to further contain, in mass %, one or more components of the following groups a to c, wherein,

Group a: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group b: 1 or more of Ti and V, the total amount is less than 0.1%;

Group c: 1 or 2 kinds of Ca and REM, the total amount is 0.0010-0.010%.

The seventh invention is a high-strength cold-rolled steel sheet with a high r value, excellent strain age hardening properties and non-aging properties at room temperature, characterized in that it contains C: 0.025-0.15%, Si: 1.0% or less, Mn: 2.0% or less, P: 0.08% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050 to 0.0250%, and N/Al is 0.3 or more, containing 0.0010% or more of solid solution state of N, the remainder is composed of Fe and unavoidable impurities, and contains an average grain size of 80% or more in area ratio: a ferrite phase of 10 μm or less, and a second phase of The structure of the martensite phase of 2% or more represented by the area ratio has an r value of 1.2 or more.

In the seventh aspect of the present invention, in addition to the aforementioned composition, it is preferable to further contain one or more components of the following groups d to g expressed in mass %, wherein,

Group d: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group e: one or more of Nb, Ti, V, the total amount is less than 0.1%;

Group f: B below 0.0030%;

Group g: one or two of Ca and REM, the total amount is 0.0010-0.010%.

The eighth present invention is a method for producing a high-strength cold-rolled steel sheet having a high r-value of 1.2 or more, excellent strain-aging hardening characteristics, and non-aging properties at room temperature, characterized in that: expressed in mass %, containing C: 0.025% to 0.15%, Si: 1.0% or less, Mn: 2.0% or less, P: 0.08% or less, S: 0.02% or less, Al: 0.02% or less, N: 0.0050 to 0.0250%, and N/Al is 0.3 The steel slab with the above composition is heated to the slab heating temperature: above 1000°C, rough rolling is performed to form a thin slab, and the thin slab is sequentially subjected to finish rolling with a discharge temperature of above 800°C, and the coiling temperature is: The hot rolling process of coiling below 650°C to form a hot-rolled sheet, the cold-rolling process of pickling and cold-rolling the hot-rolled sheet to form a cold-rolled sheet, the annealing temperature of the cold-rolled sheet: above the recrystallization temperature ~ Carry out box annealing below 800°C, then perform continuous annealing at annealing temperature: Ac ₁ transformation point ~ (Ac ₃ transformation point - 20°C), and then cool to below 500°C at a cooling rate of 10 to 300°C/sec Cold-rolled sheet annealing process in the temperature zone.

In the eighth invention, it is preferable to carry out the overaging treatment with a residence time of 20 seconds or more in a temperature range of not lower than the cooling stop temperature of the aforementioned cooling but not lower than 350° C. following the cooling after the aforementioned continuous annealing.

Regarding the eighth invention, in addition to the aforementioned composition, it is preferable to further contain one or more components of the following groups d to g expressed in mass %, wherein,

Group d: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group e: one or more of Nb, Ti, V, the total amount is less than 0.1%;

Group f: B below 0.0030%;

Group g: 1 or 2 kinds of Ca and REM, the total amount is 0.0010-0.010%.

A brief description of the drawings

Fig. 1 is a graph showing the relationship between the steel composition (N%-14/93·Nb%-14/27·Al%-14/11·B%) and the increase in tensile strength (ΔTS) after forming.

Fig. 2 is a graph showing the relationship between the B content and ΔTS in steel to which Nb and B are added together.

Fig. 3 is a graph comparing and showing the difference in the amount of increase in tensile strength due to post-forming heat treatment in a low-temperature region between steel B (conventional steel) with a large amount of solid-solution C and steel A (invention steel) with a large amount of solid-solution N picture.

Fig. 4 is a graph showing the decrease in elongation (ΔE1 ) and the effect of the increase in tensile strength (ΔTS) after forming.

FIG. 5 is a graph showing the relationship between TS×r value, BH, ΔTS, and N/(Al+Nb+B).

FIG. 6 is a graph showing the relationship between TS×r value, BH, ΔTS, and B amount.

Best way to implement the invention

In the first invention, the reason for limiting the component composition of the steel sheet to the aforementioned range will be described.

C: less than 0.01% by mass

The smaller the amount of C is, the better the deep drawability is, which is advantageous in terms of press formability. In addition, during the annealing process after cold rolling, NbC redissolves, and the solid solution C in the crystal grains increases, which tends to lower the aging resistance at room temperature. Therefore, it is preferable to suppress the amount of C to less than 0.01% by mass. More preferably, it is 0.0050 mass % or less, and especially preferably 0.0030 mass % or less.

Si: 0.005 to 1.0% by mass

Si is an effective component for suppressing the reduction of elongation and improving the strength, but when the content is less than 0.005% by mass, the effect of addition is insufficient. On the other hand, if it exceeds 1.0% by mass, the surface properties will deteriorate and the ductility will decrease. Therefore, Si It is limited to the range of 0.005 to 1.0% by mass. More preferably, it is in the range of 0.01 to 0.75% by mass.

Mn: 0.01 to 1.5% by mass

Mn is not only useful as a strengthening component of steel, but also forms MnS, which has the effect of suppressing embrittlement caused by S, but when the content is less than 0.01% by mass, the addition effect is insufficient. On the other hand, when it exceeds 1.5% by mass, it causes Deterioration of surface properties and decrease in ductility, Mn is specified to be contained in the range of 0.01 to 1.5% by mass. More preferably, it is 0.10 to 0.75% by mass.

P: 0.10% by mass or less

P effectively contributes to the strengthening of steel as a solid-solution strengthening component, but if the added amount exceeds 0.10% by mass, phosphides such as (FeNb)xP are formed, thereby reducing deep drawability. Therefore, P is limited to 0.10% by mass or less.

S: 0.01% by mass or less

When a large amount of S is contained, the amount of inclusions increases and the ductility decreases. Therefore, it is desirable to avoid the incorporation of S as much as possible, but it can be tolerated up to 0.01% by mass.

Al: 0.005 to 0.030% by mass

Al is added as a deoxidizer to improve the effective utilization of carbonitride-forming components. When the content is less than 0.005% by mass, there is no sufficient effect. As the amount of added N increases, slab defects are likely to occur during steelmaking. Therefore, Al is specified to be contained in the range of 0.005 to 0.030% by mass.

N: 0.005 to 0.040% by mass

In the present invention, N is an important element that functions to impart strain-age hardening properties to the steel sheet. However, when the content is less than 0.005% by mass, sufficient strain age hardening properties cannot be obtained. On the other hand, when the amount is increased to more than 0.040% by mass, the press formability will be reduced. Therefore, N is specified to be contained in the range of 0.005 to 0.040% by mass. Furthermore, it is preferably 0.008 to 0.015% by mass.

B: 0.0001 to 0.003% by mass

B is added in combination with Nb to effectively refine the hot-rolled structure and the cold-rolled recrystallized structure, and also to improve the resistance to secondary processing embrittlement. However, if the content is less than 0.0001% by mass, a sufficient refining effect cannot be obtained. On the other hand, if the content exceeds 0.003% by mass, not only the amount of BN precipitation increases, but also the melting zone in the slab heating stage come to obstacles. Therefore, B is contained in the range of 0.0001 to 0.003% by mass. Furthermore, it is preferably 0.0001 to 0.0015% by mass, more preferably 0.0007 to 0.0012% by mass.

Nb: 0.005 to 0.050% by mass

The compound addition of Nb with B effectively contributes to the refinement of hot-rolled structure and cold-rolled recrystallization annealed structure, and also has the effect of fixing solid solution C in the form of NbC. In addition, Nb forms a nitride called NbN, which contributes to the refinement of the cold rolling recrystallization annealing structure. However, when Nb is less than 0.005% by mass, it is not only difficult to precipitate and fix solid solution C, but also the refinement of the hot-rolled structure and the cold-rolled recrystallization annealed structure becomes insufficient. On the other hand, when it exceeds 0.050% by mass, the Reduced ductility. Therefore, Nb is contained in the range of 0.005 to 0.050% by mass. Preferably, it is 0.010 to 0.030% by mass.

In addition, as described above, Nb has the function of immobilizing solid-solution C in the form of NbC. Also, a nitride called NbN is formed. Likewise, Al and B form AlN and BN, respectively. Therefore, it is important to satisfy the relationship of the following formulas (1) and (2) in order to sufficiently secure the amount of solid-solution N and at the same time sufficiently reduce the amount of solid-solution C.

N%≥0.0015+14/93 Nb%+14/27 Al%+14/11 B% ---(1)

C%≤0.5·(12/93)·Nb% ---(2)

In addition, in this invention, in order to obtain high strain aging characteristics while preventing aging deterioration, it is suitable to reduce the crystal grain size.

That is, as shown in Fig. 4, by reducing the grain size to 20 μm or less, the mass % and containing a large amount of solid solution N, it is also possible to suppress ΔE1 to 2.0% or less. Furthermore, it is more desirable to reduce the crystal grain size d to 15 μm or less. This is because, as shown in FIG. 4 , when the crystal grain size d is reduced to 15 μm or less, it becomes possible to suppress ΔE1 to 1.5% or less.

The production conditions of the second invention will be described.

The steel with the above-mentioned suitable composition is smelted by a known smelting method such as a converter, and is made into a billet by ingot casting or continuous casting.

Next, after heating and soaking the steel slab, it is hot-rolled to form a hot-rolled sheet. In this invention, the heating temperature of hot rolling is not particularly specified, but in order to improve deep drawability, it is advantageous to fix solid solution C and precipitate it in the form of carbides. Therefore, the heating temperature of hot rolling is 1300°C or lower. as well. In addition, in order to further improve workability, it is preferable to set it at 1150° C. or lower. However, when the heating temperature is lower than 900°C, the improvement of the workability is saturated, conversely, the rolling load during hot rolling increases, and the risk of rolling accidents increases, so the lower limit of the heating temperature is preferably 900°C.

Next, the total rolling reduction during hot rolling is preferably 70% or more. This is because if the total rolling reduction is less than 70%, the grain refinement of the hot-rolled sheet becomes insufficient.

In addition, the finish rolling in hot rolling is preferably completed in the temperature range of 960-650°C, and the hot rolling processing temperature can be in the γ region above the Ar ₃ transformation point, or in the α region below the Ar ₃ transformation point. When the hot-rolling temperature exceeds 960° C., the crystal grains of the hot-rolled sheet are coarsened, and the deep-drawability after cold rolling and annealing deteriorates. On the other hand, if the temperature is lower than 650°C, the deformation resistance increases, so that the hot rolling load increases and rolling becomes difficult.

After the above-mentioned hot finish rolling is completed, the cooling is started immediately, so it is hoped that while preventing the normal grain growth, the precipitation of AlN during the cooling process is also suppressed.

Here, the above-mentioned cooling treatment conditions are not particularly limited, but the cooling start time is preferably within 1.5 seconds, more preferably within 1.0 seconds, and most preferably within 0.5 seconds after finish rolling. This is because: when cooling immediately after rolling, the degree of undercooling becomes larger in the state of strain accumulation, so more ferrite nuclei are formed, and while promoting ferrite transformation, it suppresses the γ-phase The solid solution N diffuses into the ferrite grains, and the amount of solid solution N present on the ferrite grain boundaries increases.

In addition, the cooling rate is preferably 10° C./sec or more in order to ensure solid solution N. Furthermore, especially when the hot rolling processing temperature is above the Ar ₃ transformation point, it is more suitable to set the cooling rate at 50°C/sec in order to ensure solid solution N.

Next, the hot-rolled sheet is wound into coils. The higher the winding temperature, the more favorable the coarsening of carbides, but when it exceeds 800°C, the scales formed on the surface of the hot-rolled sheet become thicker, not only the load of scale removal increases, but also the nitrogen On the other hand, if the coiling temperature is lower than 400°C, the coiling operation will become difficult, so the coiling temperature of the hot-rolled sheet must be set at 800-400°C range.

Next, cold rolling is performed on the hot-rolled sheet, but the reduction rate in this cold rolling must be 60 to 95%. This is because when the rolling reduction in cold rolling is less than 60%, a high r value cannot be expected, and on the other hand, when it exceeds 95%, the r value decreases on the contrary.

The cold-rolled sheet subjected to the above-mentioned cold rolling is then subjected to recrystallization annealing. The annealing method may be continuous annealing or batch annealing, either of which may be used, but continuous annealing is advantageous. In addition, this continuous annealing may be either processing on a normal continuous annealing line or processing on a continuous hot-dip galvanizing line.

In addition, the annealing conditions are preferably at least 650°C and at least 5 seconds. This is because if the annealing temperature is less than 650° C. and the annealing conditions are less than 5 seconds, recrystallization cannot be completed, and thus the deep drawability decreases. In order to further improve the deep drawability, it is desirable to anneal in the ferrite single-phase region at 800° C. or higher for 5 seconds or longer.

In addition, through annealing at a higher temperature in the α+γ two-phase region, the α→γ transformation partially occurs, whereby the {111} aggregate structure develops, and the r value increases, but when the α→γ transformation is completely carried out, the aggregation The structure becomes random, so the r-value decreases and the deep-drawability is impaired.

Furthermore, the upper limit of the annealing temperature is preferably 900°C. This is because when the annealing temperature exceeds 900°C, the redissolution of carbides proceeds, the solid solution C increases excessively, and the slow aging performance decreases. In addition, when the α→γ transformation occurs, the aggregate structure becomes random, so the r value Reduced, deep drawing ductility is impaired.

In addition, in the temperature rise process of the above-mentioned recrystallization annealing, the temperature range from 500° C. to the recrystallization temperature is gradually heated to sufficiently precipitate AlN, thereby effectively reducing the grain size of the steel sheet.

Here, the temperature range in which the above-mentioned controlled heating should be performed is from 500° C. at which AlN and the like start to precipitate to the recrystallization temperature.

In addition, the rate of temperature increase is preferably 1 to 20°C/sec. This is because if the temperature increase rate exceeds 20°C/sec, a sufficient amount of precipitation cannot be obtained, and on the other hand, if it is less than 1°C/sec, the precipitates are coarsened and the effect of inhibiting grain growth is weak.

In addition, after the above-mentioned recrystallization annealing, in order to further perform shape correction and surface roughness adjustment, temper rolling of 10% or less may be performed.

In addition, the cooling rate after soaking in the recrystallization annealing is preferably 10 to 50°C/sec. This is because, when the cooling rate is less than 10° C./sec, the crystal grains grow during cooling, the crystal grains become coarser, and the strain aging characteristics and the aging characteristics at room temperature decrease. On the other hand, at 50° C./sec or more, solid-solution N does not sufficiently diffuse to grain boundaries, and the aging characteristics at room temperature deteriorate. Furthermore, it is preferably 10 to 30°C/sec.

Following the above-mentioned recrystallization annealing, if necessary, a hot-dip galvanizing treatment is performed, followed by a heating alloying treatment, whereby an alloyed hot-dip galvanized steel sheet is formed.

Such hot-dip galvanizing treatment and alloying treatment are not particularly limited, and may be performed according to conventionally known methods.

Furthermore, after the alloyed hot-dip galvanized steel sheet is formed, the temper-rolled steel sheet (passive-treated steel sheet, bright-finished steel sheet, and surface with a specific roughness pattern) is applied to improve workability and appearance after processing. Steel plate), steel plate with anti-rust oil, lubricating oil and other oil film layers on the surface, etc., which are generally used as thin steel plates and have been subjected to surface treatment, if they fall within the scope of the invention, the effect of the invention can be fully obtained.

In this way, a cold-rolled steel sheet and an alloyed hot-dip galvanized steel sheet can be obtained that not only have excellent deep drawability, but also have excellent strain age hardening characteristics in which tensile strength is increased by press forming-heat treatment.

The reason for limiting the component composition of the steel sheet to the aforementioned range in the third invention will be described.

C: less than 0.01% by mass

The smaller the amount of C is, the better the deep drawability is, which is advantageous in terms of press formability. In addition, during the annealing process after cold rolling, re-dissolution of NbC proceeds, and the solid solution C in the crystal grains increases, which easily leads to a decrease in room temperature aging resistance. Therefore, it is preferable to suppress the amount of C to less than 0.01% by mass. More preferably 0.0050% by mass or less, most preferably 0.0030% by mass or less. Furthermore, from the viewpoint of ensuring strength and preventing grain coarsening, it is desirable that C is contained in an amount of 0.0005% or more. Si: 0.005 to 1.0% by mass

Si is an effective component that suppresses the decrease in elongation and improves strength, but when the content is less than 0.005% by mass, the effect of addition is insufficient. On the other hand, if it exceeds 1.0% by mass, the surface properties deteriorate and the ductility decreases, so Si is limited. It is in the range of 0.005 to 1.0% by mass. More preferably, it is the range of 0.01-0.75 mass %. Mn: 0.01 to 1.5% by mass

Mn is not only useful as a strengthening component of steel, but also forms MnS, which has the effect of suppressing embrittlement caused by S, but when the content is less than 0.01% by mass, the addition effect is insufficient. On the other hand, when it exceeds 1.5% by mass, Mn causes deterioration of surface properties and reduction in ductility, so Mn is contained in the range of 0.01 to 1.5% by mass. More preferably, it is 0.10 to 0.75% by mass. P: 0.10% by mass or less

P effectively contributes to the strengthening of steel as a solid solution strengthening component, but if the added amount exceeds 0.10% by mass, phosphides such as (FeNb)xP are formed, thereby reducing deep drawability. Therefore, P is limited to 0.10% by mass or less. S: 0.01% by mass or less

When a large amount of S is contained, the amount of inclusions increases and the ductility decreases. Therefore, it is desirable to avoid the incorporation of S as much as possible, but it is allowable up to 0.01% by mass. Al: 0.005 to 0.030% by mass

As a deoxidizer, Al is also added to improve the effective utilization of carbonitride-forming components, but if the content is less than 0.005% by mass, there is no sufficient effect. On the other hand, when the amount exceeds 0.030% by mass, it will lead to a The increased amount of N added tends to cause slab defects during steelmaking. Therefore, Al is contained in the range of 0.005 to 0.030% by mass. N: 0.005 to 0.040% by mass

In the present invention, N is an important element that functions to impart strain age hardening properties to the steel sheet. However, if the content is less than 0.005% by mass, sufficient strain age hardening properties cannot be obtained. On the other hand, when added in a large amount exceeding 0.04% by mass, the press formability will be reduced. Therefore, N is contained in the range of 0.005 to 0.040% by mass. Furthermore, it is preferably 0.008 to 0.015% by mass. B: 0.0001 to 0.003% by mass

B is added in combination with Nb to effectively refine the hot-rolled structure and the cold-rolled recrystallized structure, and also to improve the resistance to secondary processing embrittlement. However, if the content is less than 0.0001% by mass, a sufficient refining effect cannot be obtained. On the other hand, if the content exceeds 0.003% by mass, not only the amount of BN precipitation increases, but also the solid solution of the slab in the heating stage is hindered. . Therefore, B is contained in the range of 0.0001 to 0.003% by mass. Furthermore, it is preferably 0.0001 to 0.0015% by mass, more preferably 0.0007 to 0.0012% by mass. Nb: 0.005-0.050%, Ti: 0.005-0.070%, V: 0.005-0.10%

Nb, Ti, and V are added in combination with B, which helps to refine the hot-rolled structure and cold-rolled recrystallized structure, and has the effect of making C precipitate in the form of NbC, TiC, and VC, so they are added together with B as needed, but When each is less than 0.005 mass, the effect is insufficient. On the other hand, when Nb exceeds 0.05%, Ti exceeds 0.070%, and V exceeds 0.10%, ductility deteriorates. Therefore, Nb is 0.005-0.050%, Ti is 0.005-0.070%, and V is 0.005-0.10%.

In addition, as described above, Nb has the function of fixing solid-solution C in the form of NbC. A nitride called NbN is also formed. Likewise, Al and B form AlN and BN, respectively. Therefore, it is important to satisfy the relationship of the following formulas (1) and (2) in order to sufficiently reduce the amount of solid solution N while sufficiently securing the amount of solid solution N.

N%≥0.0015+14/93 Nb%+14/27 Al%+14/11 B% ---(1)

C%≤0.5 (12/93) Nb% ---(2) N/Al or N/(Al+Nb+Ti+V+B): 0.30 or more

Al forms AlN to reduce solid solution N. In order to secure an appropriate amount of solid solution N, N/Al must be 0.30 or more. In addition, when adding Nb, Ti, V or B in combination, they all form NbN, TiN, VN, and BN respectively to reduce solid solution N, so in order to ensure an appropriate amount of solid solution N, it is necessary to make N/(Al+Nb+Ti+ V+B) is 0.30 or more. Solid solution N: 0.0010% or more

In order to improve the strain age hardening properties of the steel sheet, solid solution N must be present in a content of 0.0010% or more.

Here, the amount of solid solution N is obtained by subtracting the amount of precipitated N from the total amount of N in the steel. In addition, the present inventors compared and studied various analysis methods as an analysis method for the amount of precipitated N, and based on the results, it is effective to obtain it by an electrolytic extraction analysis method using a constant potential electrolysis method. In addition, as a method of dissolving matrix iron for extraction analysis, there are acid decomposition method, halogen method, and electrolysis method. Among them, the electrolysis method does not decompose extremely unstable precipitates such as carbides and nitrides, and can only dissolve matrix iron stably. An acetylacetone system was used as the electrolytic solution, and it was electrolyzed at a constant potential. In the present invention, the result of measuring the amount of precipitated N by the constant potential electrolysis method shows the best correspondence with the actual component strength.

Therefore, in the present invention, the residue extracted by the constant potential electrolysis method is chemically analyzed to determine the amount of N in the residue, which is taken as the amount of precipitated N.

Furthermore, in order to obtain higher BH and ΔTS, the amount of solid solution N is preferably 0.0015% or more, more preferably 0.0020% or more, particularly preferably 0.0030% or more.

The cold-rolled steel sheet of the present invention is a cold-rolled steel sheet for deep drawing that has the above-mentioned composition and is excellent in strain age hardening characteristics characterized by a TS×r value ≥ 750 MPa.

A steel plate with a TS×r value of less than 750 MPa cannot be widely applied to members having elements of structural members. In addition, in order to further expand the applicable range, the TS×r value is preferably 850 MPa or more.

Conventional painting and baking conditions have been adopted as a standard of 170°C x 20 minutes. Furthermore, when a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solid solution N, hardening can be achieved even with a milder (lower temperature) treatment. In other words, aging conditions can be obtained in a wider range. . In addition, in general, in order to strive for the amount of hardening, it is advantageous to keep at a higher temperature for a longer time as long as it does not soften due to excessive aging.

Specifically, for the steel sheet of the present invention, the lower limit of the heating temperature at which hardening after pre-deformation becomes significant is about 100°C. On the other hand, when the heating temperature exceeds 300° C., the hardening reaches its peak, and conversely, a slight softening tendency occurs, and the generation of thermal strain and temper color becomes remarkable. In addition, regarding the holding time, when the heating temperature is about 200° C., if it is about 30 seconds or longer, sufficient curing can basically be achieved. In order to further obtain large and stable hardening, the holding time is preferably 60 seconds or more. However, if the holding time exceeds 20 minutes, further hardening cannot be expected, but production efficiency will be significantly lowered, which is disadvantageous in practical use.

In view of the above, in the present invention, as the aging treatment conditions, the heating temperature of 170° C., which is the conventional paint baking treatment condition, was evaluated for 20 minutes, and the holding time was set. The steel sheet of the present invention can stably achieve sufficient hardening even under the aging treatment conditions of low-temperature heating and short-time holding that conventional paint-bake type steel sheets cannot achieve sufficient hardening. Furthermore, the method of heating is not particularly limited. In addition to the atmosphere heating carried out by the furnace used in the usual paint baking, for example, heating by induction heating, non-oxidizing flame, laser, plasma, etc. can be carried out very well. ground use. In addition, only the portion where the strength is to be increased may be selectively heated.

The strength of automotive parts must be able to resist complex stress loads from the outside. Therefore, for the raw material steel plate, not only the strength characteristics in the small strain region but also the strength characteristics in the large strain region are important. Based on this point, the present inventors made BH of the steel sheet of the present invention to be 80 MPa or more and ΔTS of 40 MPa or more to be used as a material for automobile parts. Furthermore, it is more preferable that BH is 100 MPa or more, and ΔTS is 50 MPa or more. In order to increase BH and ΔTS, it is only necessary to set the heating temperature during the aging treatment to a higher temperature side and/or set the holding time to a longer time side.

In addition, the steel sheet of the present invention has the conventional advantage that aging deterioration (a phenomenon in which YS increases and E1 decreases) does not occur even if it is left at room temperature for a long time of about one year without being formed.

In addition, in the present invention, even if hot-dip galvanizing or alloying hot-dip galvanizing is performed on the surface of the above-mentioned cold-rolled steel sheet of the present invention, there is no problem, and TS, BH, and ΔTS are shown at the same level as before plating. In addition, electrolytic zinc plating, electrolytic tin plating, electrolytic chromium plating, electrolytic nickel plating, and the like can be suitably used as types of plating other than hot-dip galvanizing and alloying hot-dip galvanizing.

The production conditions of the fourth invention will be described.

First, a steel with a composition containing C: less than 0.01%, N: 0.0050 to 0.04%, Al: 0.005 to 0.03%, Si: 0.005 to 1.0%, and Mn: 0.01 to 1.5% is smelted using a conventionally known smelting method such as a converter. %, P: 0.1% or less, S: 0.01% or less, or one of Nb: 0.005-0.050%, Ti: 0.005-0.070%, V: 0.005-0.10% together with B: 0.0001-0.003%, or 2 or more types, and N/(Al+Nb+Ti+V+B): 0.30 or more. Then, it is solidified into a steel ingot by ingot casting method or continuous casting method.

The steel ingot is heated and soaked, and then hot-rolled to form a hot-rolled sheet. When the heating temperature (SRT) is too low, the workability improvement effect is saturated, and the rolling load during hot rolling increases, and there is a danger of rolling accidents or insufficient homogenization of solid solution N, so SRT is 950 °C or higher is better. Furthermore, in order to improve the deep drawability, it is advantageous to fix solid solution C to precipitate in the form of carbides, and the SRT is preferably below 1300°C. In addition, in order to further improve processability, it is preferable to set it as 1150 degreeC or less.

If the total rolling reduction from the rough rolling to the finish rolling of the hot rolling is less than 80%, the grain refinement of the hot-rolled sheet is insufficient, so it is better to make it 80% or more.

When the rough rolling temperature exceeds 1000°C, the grains of the γ→α phase transition will be coarsened, and the r value will decrease. When it is less than Ar ₃ , the r value will decrease due to the recrystallization of the α grain or the grain growth, so Rough rolling is preferably carried out in a temperature range of 1000°C or lower and Ar ₃ or higher.

On the other hand, when the finish rolling is completed in a temperature range exceeding Ar ₃ , due to the γ→α transformation, the aggregate structure becomes random, and excellent deep drawability cannot be obtained. On the other hand, it ends at a temperature lower than 600°C. In finish rolling, no further improvement in deep drawability is expected, but only an increase in rolling load, so finish rolling is preferably performed in a temperature range of Ar ₃ or lower and 600°C or higher.

In addition, if lubricated rolling is not performed during finish rolling, due to the frictional force between the roll and the steel plate, an additional shearing force acts on the surface of the steel plate. Ideal {110} orientation, so deep drawability deteriorates. Therefore, it is better to perform finish rolling while lubricating.

Next, the hot-rolled sheet is wound into a coil. In addition, the processed material which passed through the winding process is also called a coil. The higher the coiling temperature (CT) of the hot-rolled sheet is, the more favorable it is for the coarsening of carbides, but when it exceeds 800°C, the scales formed on the surface of the hot-rolled sheet become thicker, and the load of removing the scales increases. , or nitrides are formed, resulting in fluctuations in the amount of solid solution N in the longitudinal direction of the coil. On the other hand, if the temperature is lower than 400°C, the winding operation becomes difficult. For this reason, it is better to take CT as 800-400°C.

Next, the obtained hot-rolled sheet is subjected to recrystallization annealing by continuous annealing or batch annealing. This annealing (hot-rolled sheet annealing) is performed to recrystallize the rolled texture formed by the α-zone intermediate temperature rolling performed in the finish rolling, thereby obtaining a recrystallized texture.

Next, the hot-rolled sheet is cold-rolled to become a cold-rolled sheet. When the reduction ratio of cold rolling is less than 60%, a high r value cannot be expected. On the other hand, if it exceeds 95%, the r value will decrease instead, so it is preferably 60 to 95%.

Next, the cold-rolled sheet is recrystallized annealed. This annealing is preferably performed on either a continuous annealing line or a continuous hot-dip galvanizing line. The annealing condition is preferably an annealing temperature of 650° C. or higher and a holding time of 5 seconds or more. If any one of the annealing temperature of 650° C. or higher and the holding time of 5 seconds or longer is not satisfied, recrystallization cannot be completed, and deep-drawability decreases. Furthermore, in order to obtain more excellent deep drawability, it is preferable that the annealing temperature is 800° C. or higher and the holding time is 5 seconds or higher. However, when the annealing temperature exceeds 900°C, the redissolution of carbides proceeds, and the solid solution C increases excessively, so the slow aging performance (resistance to aging at room temperature) decreases, and when the α→γ transformation occurs, the aggregate structure becomes random. , the r value decreases, and the deep drawability is damaged, so the annealing temperature is preferably below 900 °C.

Furthermore, a cold-rolled annealed sheet obtained by recrystallizing and annealing a cold-rolled steel sheet is subjected to hot-dip galvanizing or re-alloying treatment as necessary. At this time, in the plating treatment, the cooling rate from after recrystallization annealing to before the plating treatment is 5°C/sec or more, and the plate temperature during hot-dip galvanizing is preferably 400-600°C. , it is better to set the treatment temperature at 400-600° C. and the treatment time at 5-40 seconds.

Furthermore, the cold-rolled steel sheet or hot-dip galvanized steel sheet after recrystallization annealing may be temper-rolled for shape correction and surface smoothness adjustment. The reduction ratio of the temper rolling is preferably 10% or less. This is because the r value decreases when the reduction ratio exceeds 10%.

The reasons for limiting the composition of the high-strength cold-rolled steel sheet of the fifth invention will be described.

C: 0.0015～0.025%

C controls the structure to be uniform and fine, and in order to secure a sufficient amount of acicular ferrite, it must be contained in an amount of 0.0015% or more in the present invention. On the other hand, when it exceeds 0.025%, the carbide fraction in the steel sheet is too large, and the ductility, r-value, and formability are remarkably lowered. In this case, C is limited in the range of 0.0015 to 0.025%. Furthermore, from the viewpoint of improving formability, it is preferably 0.020% or less, more preferably 0.010% or less. In addition, especially from the standpoint of stabilizing the BH amount and material, it is more preferable that the C content exceeds (12/93) Nb (%) (here, Nb is the Nb content (%)).

Si: 1.0% or less

Si is a useful element that can increase the strength of the steel sheet without significantly reducing the ductility of the steel. In the present invention, it is preferably contained at least 0.005%, and when high strength is particularly required, it is more preferably contained at least 0.10%. On the other hand, Si greatly increases the transformation point during hot rolling, making it difficult to ensure the shape, or also adversely affects the surface properties, surface chemical treatment properties, etc., especially the aesthetics of the steel sheet surface, and further , have a bad effect on the coating property. In the present invention, it is limited to 1.0% or less. If Si is 1.0% or less, the above-mentioned bad influence can be suppressed very low. In addition, for applications that particularly require the appearance of the surface of the plated steel sheet, Si is desirably contained at 0.5% or less.

Mn: 2.0% or less

Mn is an effective element for preventing hot cracks caused by S, and it is better to add according to the amount of S contained. In addition, Mn has a great effect on the refinement of crystal grains, and it is desirable to add Mn for improving the material. From the viewpoint of stably fixing S, it is desirable that Mn be contained in an amount of 0.1% or more. In addition, Mn is an element that increases the strength of the steel sheet, and when more strength is required, it is desirable to contain 0.5% or more. Furthermore, it is more preferably 0.8% or more.

When the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet, especially the dispersion of the strain age hardening characteristics against fluctuations in hot rolling conditions are remarkably improved. However, when the excessive content of Mn exceeds 2.0%, although the detailed mechanism is unclear, it tends to increase the hot deformation resistance, and also tends to deteriorate the weldability and weld formability, thereby significantly suppressing the formation of ferrite. The ductility is remarkably lowered, and the r-value also tends to be significantly lowered, so Mn is limited to 2.0% or less. For applications requiring better corrosion resistance and formability, it is preferably 1.5% or less.

P: less than 0.1%

P is an element useful as a solid-solution strengthening element of steel, and from the viewpoint of increasing strength, it is preferable to contain 0.002% or more, especially when high strength is required, it is preferable to contain 0.02% or more. On the other hand, when it is contained excessively, the steel is embrittled, and the hemming workability of the steel sheet is further deteriorated. In addition, since P has a strong tendency to segregate in steel, it causes embrittlement of the weld zone due to it. Therefore, P is limited to 0.1% or less. In addition, in applications where the hemming workability and weld toughness are particularly important, P is preferably 0.08% or less. More preferably 0.06% or less.

S: 0.02% or less

S is an element that exists in the form of inclusions in the steel sheet, reduces the ductility of the steel sheet, and also causes corrosion resistance to deteriorate. It is better to reduce it as much as possible. In the present invention, S is limited to 0.02% or less. Especially for applications requiring good workability, S is preferably 0.015% or less. In addition, when particularly excellent hemming workability is required, S is preferably 0.010% or less. Also, although the detailed mechanism is unclear, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening characteristics of the steel sheet at a high level.

Al: less than 0.02%

Al is an element that functions as a deoxidizer, improves the purity of steel, and refines the structure of the steel sheet. In the present invention, it is desirable to contain 0.001% or more. In the present invention, solid-solution N is used as a strengthening element, but the aluminum-killed steel containing Al in an appropriate range has better mechanical properties than the conventional ebullient steel without adding Al. On the other hand, excessive Al content deteriorates the surface properties of the steel sheet and significantly reduces solid-solution N, making it difficult to obtain a very large amount of strain age hardening that is the focus of the present invention. In this case, in the present invention, Al is limited to 0.02% or less. Furthermore, from the viewpoint of the stability of the material, it is more desirable to make Al 0.001 to 0.015%. In addition, there is also concern about grain coarsening when the Al content is reduced, but in the present invention, this is effectively prevented by setting the optimum amount of other alloy elements and setting the annealing conditions within the optimum range.

N: 0.0050～0.0250%

N is an element that increases the strength of a steel sheet through solid solution strengthening and strain age hardening, and is the most important element in the present invention. In addition, in the present invention, by containing an appropriate amount of N, the Al content is adjusted to an appropriate value as described above, and manufacturing conditions such as hot rolling conditions and annealing conditions are also controlled to ensure necessary and sufficient Al content in cold rolled products or plated products. of solid solution N. Accordingly, the effect of increasing the strength (yield stress and tensile strength) by solid solution strengthening and strain age hardening can be fully exerted, and the following target values of the mechanical properties of the steel sheet of the present invention can be stably obtained: tensile strength 340 MPa or more, baked The amount of bake hardening (BH amount) is 80 MPa or more, and the increase in tensile strength ΔTS before and after strain aging treatment is 40 MPa or more. In addition, N has the effect of lowering the transformation point, and its inclusion is effective in the case of rolling a thin product that does not want to be rolled, which greatly lowers the transformation point.

When N is less than 0.0050%, it is difficult to stably express the aforementioned effect of increasing the strength. On the other hand, when N exceeds 0.0250%, the occurrence rate of internal defects in the steel sheet increases, and slab cracks during continuous casting often occur. Therefore, N is limited to the range of 0.0050 to 0.0250%. Furthermore, N is preferably 0.0070 to 0.0200%, more preferably 0.0100 to 0.0170%, from the viewpoint of material stability in the overall manufacturing process and yield improvement. In addition, as long as the amount of N is within the range of the present invention, there is no adverse effect on weldability or the like at all.

Solid solution N: 0.0010% or more

In cold-rolled products, in order to ensure sufficient strength and effectively exert strain age hardening caused by N, at least 0.0010% or more of solid-solution N (also called solid-solution N) must exist in the steel sheet.

Here, the amount of solid solution N is the amount of precipitated N subtracted from the total amount of N in the steel, and the obtained value is taken as the amount of solid solution N. Furthermore, the present inventors compared and studied various methods as an analysis method for the amount of precipitated N, and as a result, it is effective to obtain it by an electrolytic extraction method using a constant potential electrolysis method. In addition, as a method of dissolving matrix iron for extraction analysis, there are acid decomposition method, halogen method, and electrolysis method. Among them, the electrolytic method does not decompose extremely unstable precipitates such as carbides and nitrides, and can only dissolve matrix iron stably. An acetylacetone system was used as the electrolytic solution, and it was electrolyzed at a constant potential. In the present invention, the results of measuring the amount of precipitated N using the constant potential electrolysis method showed a good correspondence with actual material changes.

In view of this, in the present invention, the residue extracted by the constant potential electrolysis method is chemically analyzed to determine the amount of N in the residue, which is taken as the amount of precipitated N.

Furthermore, when a higher amount of BH and ΔTS are required, it is preferable to make the amount of solid solution N 0.0020% or more, and to obtain a higher value, it is preferable to make it 0.0030% or more. The upper limit of the solid-solution N amount is not particularly limited, but the total N amount is small even if the overall residual mechanical properties are lowered.

N/Al (ratio of N content to Al content): 0.3 or more

In order to stably leave 0.0010% or more of solid-solution N in the product state, the amount of Al, which is an element that strongly fixes N, must be limited. As a result of examining steel sheets in which combinations of N content (0.0050 to 0.0250%) and Al content (0.02% or less) were widely varied within the composition range of the present invention, it was found that by setting N/Al to 0.3 or more, The solid solution N in cold-rolled products and coated products is stably maintained at 0.0010% or more. Therefore, N/Al is limited to 0.3 or more. Furthermore, from the viewpoint of stably improving the strain age hardening properties, N/Al is preferably 0.6 or more. More preferably, it is 0.8 or more.

Nb: 0.002 to 0.050%

Nb is compounded with B to effectively generate the acicular ferrite phase, and in the present invention, it must be contained in an amount of 0.002% or more. On the other hand, when the content exceeds 0.050%, the effect is saturated, and the hot deformation resistance increases remarkably, making hot rolling difficult. Therefore, Nb is limited within the range of 0.002 to 0.050%. Furthermore, it is more preferably 0.005 to 0.040%.

B: 0.0001～0.0050%

B is an element that is complexed with Nb and effectively functions to form an acicular ferrite phase, and in the present invention, it must be contained in an amount of 0.0001% or more. On the other hand, when the content exceeds 0.0050%, the solid solution N that contributes to the strain age hardening characteristic is reduced. Therefore, B is limited in the range of 0.0001 to 0.0050%. Furthermore, it is preferably 0.0003 to 0.0030%. More preferably, it is 0.0005 to 0.0030%.

In the present invention, in addition to the above composition, it is preferable to further contain one or more of the following groups a to c, wherein,

Group a: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group b: 1 or more of Ti and V, the total amount is less than 0.1%;

Group c: 1 or 2 kinds of Ca and REM, the total amount is 0.0010-0.010%.

Group a elements: Cu, Ni, Cr, and Mo are all elements that contribute to the increase in the strength of the steel sheet, and can be selected and contained alone or in combination as needed. This effect is seen when Cu, Ni, Cr, and Mo are contained in an amount of 0.01% or more. However, when the content is too large, the thermal deformation resistance increases, or the chemical surface treatability and generalized surface treatment characteristics deteriorate, and the weld zone hardens, and the weld zone formability deteriorates. Therefore, each of Cu, Ni, Cr, and Mo individually is preferably 1.0% or less, 1.0% or less, 0.5% or less, and 0.2% or less. When they are contained in combination, the total amount is preferably 1.0% or less.

Group b elements: Ti and V are elements that contribute to grain refinement and homogenization, and can be selected and contained alone or in combination as needed. This effect is seen when Ti and V are contained in an amount of 0.005% or more, respectively. However, when the content is too large, the thermal deformation resistance increases while chemical surface treatment and surface treatment properties in a broad sense deteriorate. Further, there is also an adverse effect of reducing solid solution N. Therefore, Ti and V alone are preferably 0.1% or less and 0.1% or less, respectively. When compounded, the total amount is preferably 0.1% or less.

Group c elements: Ca and REM are elements that play a role in controlling the shape of inclusions, especially when curling formability is required, it is better to contain them alone or in combination. When the total amount of group d elements is less than 0.0010%, the effect of controlling the morphology of inclusions is insufficient, while on the other hand, when it exceeds 0.010%, the occurrence of surface defects becomes significant. Therefore, it is preferable to limit the total amount of the d group elements to the range of 0.0010 to 0.010%. According to this, the hemming workability can be improved without being accompanied by the occurrence of surface defects.

The structure of the steel sheet of the present invention will be described.

The steel sheet of the present invention has a structure composed of 5% or more of acicular ferrite phase and a ferrite phase with an average crystal grain size of 20 μm or less in terms of area ratio.

Area ratio of acicular ferrite phase: 5% or more

The cold-rolled steel sheet of the present invention contains 5% or more of the acicular ferrite phase in terms of area ratio. By the presence of 5% or more acicular ferrite, good ductility and a large amount of strain age hardening can be obtained. Although the detailed mechanism is unclear, it is presumed that the presence of the acicular ferrite phase allows strain to be stored inside very efficiently during pre-straining before aging. In addition, the presence of the acicular ferrite phase is also effective in improving the aging deterioration at room temperature and achieving room temperature non-aging properties. Furthermore, in order to obtain a good strength-ductility balance and higher strength, the area ratio of the acicular ferrite phase is preferably 10% or more. Furthermore, the presence of a large amount of acicular ferrite phase exceeding 20% has the problem of lowering the r value. Therefore, the area ratio of the acicular ferrite phase is 5% or more, preferably 10% or more and 20% or less.

The acicular ferrite phase mentioned in the present invention is a low-temperature phase transformation phase that does not have carbides in the interior, which is unique to ultra-low carbon steel with the composition of the present invention. Ferrite is clearly distinguished, and is a phase that has a high internal dislocation density and is harder than the polygonal ferrite phase.

According to optical microscope observation, the following types of acicular ferrite phases are distributed alone or in combination: ① grain boundaries with irregular edges and corners; ② grain shapes that exist along grain boundaries such as precipitates ; ③Scratch-shaped grains or grain clusters (sub-grain boundaries can be seen in most of the relatively large second-phase particles), etc. They can be clearly distinguished from usual polygonal ferrites. In addition, the corroded color tone in the crystal grains is also different from martensite and bainite, and there is almost no change from ordinary polygonal ferrite, so it can also be clearly distinguished from martensite and bainite. According to the observation of the transmission electron microscope, the dislocation density near the grain boundary and/or within the grain of the acicular ferrite phase is very high, especially in the part where the dislocation density of the acicular ferrite phase of the ③ form is very high And the lower part becomes layered.

The cold-rolled steel sheet of the present invention is aimed at steel sheets for automobiles requiring high formability, and phases other than the acicular ferrite phase are ferrite phases in order to ensure ductility. When the area ratio of the ferrite phase is less than 80%, it becomes difficult to secure the ductility and high r-value required for a steel sheet for automobiles requiring workability. Furthermore, when better ductility is required, the area ratio of the ferrite phase is desirably 80% or more, more preferably 85% or more. In addition, the ferrite referred to in the present invention refers to so-called polygonal ferrite that does not remain in a strained state.

Average grain size of ferrite phase: 20 μm or less

In the present invention, as the average crystal grain size, the value calculated by the quadrature method prescribed by ASTM from the cross-sectional structure photograph and the nominal diameter obtained by the cutting method prescribed by the same ASTM (for example, refer to Umemoto et al.: heat treatment, 24(1984), 334), whichever is larger.

The cold-rolled steel sheet of the present invention ensures a predetermined amount of solid-solution N in the product stage, but according to the experiments and studies of the present inventors, even in steel sheets having the same amount of solid-solution N, the strain-age hardening characteristics sometimes vary. , it was found that one of the main factors is the grain size. In the structure of the present invention, by making the average crystal grain size at least 20 μm or less, preferably 15 μm or less, a high BH amount and ΔTS can be stably obtained. Although the detailed mechanism is unclear, it can be speculated that it is related to the segregation and precipitation of alloy elements to the grain boundaries, and the influence of processing and heat history on them.

Therefore, in order to stabilize the strain age hardening characteristics, the average crystal grain size of the ferrite phase is 20 μm or less, preferably 15 μm or less.

The cold-rolled steel sheet of the present invention having the above-mentioned composition and structure has a tensile strength (TS) of 340 MPa or more and about 590 MPa or less, a high r value of 1.2 or more, and excellent strain age hardening properties. Steel plates with TS less than 340MPa cannot be widely used in components with structural component elements. In addition, in order to further expand the applicable range, it is desirable that TS is 400 MPa or more. In addition, if the r value is less than 1.2, it cannot be applied to a wide range of press-formed parts. In addition, the preferable range of r value is 1.3 or more.

Conventional painting and baking conditions have been adopted as a standard of 170°C x 20 minutes. Furthermore, when a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solid solution N, hardening can be achieved even with a milder (low temperature side) treatment, in other words, aging conditions can be obtained in a wider range. . In addition, in general, in order to strive for the amount of hardening, it is advantageous to keep at a higher temperature for a longer time within the limit of not softening due to excessive aging.

Specifically, the lower limit of the heating temperature at which hardening becomes significant after pre-deformation of the steel sheet of the present invention is about 100°C. On the other hand, when the heating temperature exceeds 300°C, the hardening reaches its peak, and on the contrary, there is a tendency to slightly soften, and furthermore, the generation of thermal strain and temper color becomes remarkable. In addition, regarding the holding time, when the heating temperature is about 200° C., it is about 30 seconds or more, and sufficient curing can basically be achieved. In order to obtain greater stable hardening, the holding time is preferably 60 seconds or more. However, when the holding time exceeds 20 minutes, not only further hardening cannot be obtained, but also the production efficiency is significantly lowered, which is disadvantageous in practical use.

Therefore, in the present invention, as the aging treatment conditions, the heating temperature of 170° C., which is the conventional paint baking treatment condition, was evaluated for 20 minutes, and the holding time was set. Even under the aging treatment conditions of low-temperature heating and short-time holding that the conventional paint-bake type steel sheet cannot achieve sufficient hardening, the steel sheet of the present invention can stably achieve large hardening. Furthermore, the method of heating is not particularly limited. In addition to the atmosphere heating carried out by the furnace used in the usual paint baking, for example, heating by induction heating, non-oxidizing flame, laser, plasma, etc. can be carried out. Well adopted.

The strength of automotive parts must be able to resist complex stress loads from the outside, so the strength characteristics of the raw steel sheet are important not only in the small strain region, but also in the larger strain region. Based on this point, the present inventors made the BH amount (corresponding to the strength characteristic of relatively small strain region) of the steel plate of the present invention to be 80 MPa or more, and simultaneously made the ΔTS amount (corresponding to the strength characteristic of relatively large strain region) of the steel plate of the present invention to be the raw material of automobile parts. characteristics) is above 40MPa. Furthermore, it is more preferable that the amount of BH is 100 MPa or more, and ΔTS is 50 MPa or more. In addition, the amount of BH and ΔTS can be increased by setting the heating temperature during the aging treatment to a higher temperature and/or setting the holding time to a longer time.

The effect of the present invention can also be brought into play even when the product plate thickness is relatively thick, but when the product plate thickness exceeds 3.2mm, the necessary and sufficient cooling rate cannot be ensured in the cold-rolled plate annealing process, and strain aging occurs during continuous annealing, making it difficult The strain age hardening characteristic targeted for the product is obtained. Therefore, the thickness of the steel sheet of the present invention is preferably 3.2 mm or less.

In addition, in the present invention, there is no problem even if electroplating or hot-dipping is performed on the surface of the above-mentioned cold-rolled steel sheet of the present invention. These coated steel sheets also showed the same levels of TS, BH, and ΔTS as before plating. As the type of coating, electro-galvanizing, hot-dip galvanizing, alloyed hot-dip galvanizing, electro-tin plating, electro-chrome plating, electro-nickel plating, etc. can all be used well.

The manufacturing method of the steel plate of 6th this invention is demonstrated.

The steel plate of the present invention is basically manufactured by sequentially performing the following steps: heating a steel slab having the composition in the above-mentioned range, rough rolling it to form a thin slab, subjecting the thin slab to finish rolling, cooling after finishing rolling, and coiling to form a thin slab. A hot-rolling process of a hot-rolled sheet, a cold-rolling process of making a cold-rolled sheet by pickling and cold-rolling the hot-rolled sheet, and a cold-rolled sheet annealing step of continuously annealing the cold-rolled sheet.

The slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macrosegregation of components, but it may also be produced by an ingot casting method or a thin slab casting method. In addition, after manufacturing the slab, in addition to the conventional method of once cooling to room temperature and then reheating, straight rolling is carried out in which the hot slab is placed in a heating furnace without cooling and rolled, or is directly heated after being slightly heated. Energy-saving processes such as rolling direct rolling can also be used without problems. In particular, in-line rolling is one of useful techniques for effectively securing N in a solid solution state.

First, the reasons for limiting the conditions of the hot rolling process will be described.

Slab heating temperature: above 1000°C

The slab heating temperature is preferably 1000° C. or higher in order to ensure a necessary and sufficient amount of solid solution N in the initial state and satisfy the target value of the amount of solid solution N in the product. Furthermore, since the loss increases with the increase in oxidation weight, it is better to set it at 1280°C or lower.

The slab heated under the above conditions is rough rolled into a thin slab. In addition, the conditions of rough rolling do not need to be specifically defined, What is necessary is just to carry out by a conventional method. However, from the viewpoint of securing the amount of solid-solution N, it is desirable to perform the process in as short a time as possible. The thin slab is then finish-rolled to make hot-rolled plate.

Furthermore, in the present invention, between rough rolling and finish rolling, it is desirable to join front and rear adjacent thin slabs together and roll them continuously. It is preferable to use a pressure welding method, a laser welding method, an electron beam welding method, etc. as a joining means.

Continuous rolling eliminates so-called unsteady portions of rolling at the front and rear ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length and width of the coil (material to be processed). This is extremely effective in improving not only the hot-rolled steel sheet but also the cross-sectional shape and size of the cold-rolled steel sheet. In addition, tension can always be applied when cooling on the hot metal roller table after rolling, so the shape of the steel sheet can be maintained well.

In addition, in order to stably pass the front end of the coil by performing continuous rolling, lubricated rolling, which is difficult to use due to the problems of plate passability and biting property, can be used in the usual single-rolling of each thin slab. Accordingly, the rolling load can be reduced, and at the same time, the surface pressure of the roll can be reduced, and the life of the roll can be extended.

In addition, in the present invention, on the input side of the finish rolling mill between rough rolling and finish rolling, a sheet slab edge heater for heating the end of the sheet bar in the width direction and a sheet bar heater for heating the end of the sheet bar in the longitudinal direction are used. Either one or both of them, it is better to make the temperature distribution in the width direction and length direction of the thin slab uniform. According to this, material dispersion in the steel sheet can be further reduced. The thin slab edge heater and the thin slab heater are preferably induction heating.

The order of use is expected to firstly use thin slab edge heaters to compensate for temperature differences in the width direction. The amount of heating at this time is also based on the steel composition, but it is better to set the temperature distribution range in the width direction on the side of the finish rolling to about 20°C or less. The temperature difference in the length direction is then compensated by the thin slab heater. At this time, it is better to set the heating amount at the end of the length to be about 20°C higher than the temperature in the central zone.

Finishing rolling temperature: above 800°C

In order to obtain a uniform and fine hot-rolled mother plate structure, the finishing rolling temperature FDT is set to be above 800°C. When the FDT is lower than 800°C, the structure of the steel plate becomes inhomogeneous, and the processed structure remains partially, and the inhomogeneity of the structure cannot be eliminated after the cold rolling and annealing process and remains. Therefore, when a high winding temperature is used to avoid the residue of the processed structure, coarse crystal grains are generated, and the same problem occurs. In addition, when the coiling temperature is high, the amount of solid-solution N significantly decreases, so it is difficult to obtain the target tensile strength of 340 MPa or more. Due to this situation, the finishing rolling temperature FDT is taken as 800°C. In order to further improve the mechanical properties, it is desirable to set the FDT to 820° C. or higher. Furthermore, from the viewpoint of increasing the r value, the FDT is more preferably at least the Ac ₃ transformation point. In addition, the upper limit of FDT is not particularly defined, but when it is excessively high, occurrence of scale defects and the like becomes remarkable. Furthermore, FDT is preferably set to about 1000°C.

Winding temperature: below 800°C

As the coiling temperature CT decreases, the strength of the steel sheet tends to increase. In order to ensure the target tensile strength TS340MPa or more, it is better to take CT below 800°C. Furthermore, when the CT is less than 200° C., the shape of the steel sheet tends to be uneven, and there is a high risk of occurrence of defects in actual operation, showing a tendency for the uniformity of the material to decrease. Therefore, it is desirable that CT be set at 200°C or higher. Furthermore, in the case where material uniformity is more required, it is better to take CT above 300°C. More preferably, it is 350° C. or higher.

In addition, in the present invention, lubricated rolling may be performed in order to reduce the hot rolling load during finish rolling. There is an effect of making the shape and material of the hot-rolled sheet more uniform by performing lubricated rolling. Furthermore, the friction coefficient during lubricated rolling is preferably in the range of 0.25 to 0.10. In addition, the operation of hot rolling is stabilized by combining lubricated rolling and continuous rolling.

The hot-rolled sheet subjected to the above-mentioned hot-rolling step is then subjected to pickling and cold-rolling in the cold-rolling step to be a cold-rolled sheet.

The pickling conditions may be generally known conditions and are not particularly limited. Furthermore, when the scales of the hot-rolled sheet are extremely thin, cold rolling may be performed without pickling.

In addition, the cooling conditions may be generally known conditions, and are not particularly limited. Furthermore, from the viewpoint of ensuring the uniformity of the structure, the cold rolling reduction ratio is preferably 60% or more. The reasons for limiting the conditions of the cold-rolled sheet annealing step will be described below.

The cold-rolled sheet is then subjected to a cold-rolled sheet annealing process consisting of continuous annealing-cooling.

Continuous annealing temperature: the temperature in the ferrite-austenite two-phase coexistence region

The acicular ferrite phase is formed by annealing at a temperature in the ferrite-austenite two-phase coexistence region. Furthermore, the (111) aggregate structure is also well developed in the ferrite phase, so a high r value can be obtained. On the other hand, at a high temperature beyond the ferrite-austenite two-phase coexistence region to become austenite single phase, the aggregate structure of the steel plate is randomized through reverse phase transformation and phase transformation, so the r value decreases . For this reason, in the present invention, the annealing temperature of continuous annealing is limited to the temperature in the ferrite-austenite two-phase coexistence region above the recrystallization temperature. Furthermore, from the viewpoint of the stability of the r value, it is preferable to set the temperature at which the fraction of austenite is 10% or more and 50% or less.

In addition, the retention time of the continuous annealing time is preferably as short as possible from the viewpoint of production efficiency, microstructure refinement, and securing the amount of solid solution N. From the viewpoint of the stability of the work, the holding time is preferably 10 seconds or more. In addition, from the viewpoint of refining the structure and ensuring the amount of solid solution N, it is preferable to set it at 90 seconds or less. Furthermore, from the viewpoint of stabilization of the material, it is more preferably 20 seconds or more.

Cooling after continuous annealing: cooling to a temperature range below 500°C at a cooling rate of 10 to 300°C/sec

Cooling after soaking in continuous annealing is important from the viewpoint of refining the structure, forming the acicular ferrite phase, and ensuring the amount of solid-solution N. In the present invention, the supercode is continuously cooled to a temperature range below 500° C. at a cooling rate above 10° C./second. When the cooling rate is less than 10°C/sec, the required amount of acicular ferrite, uniform and fine structure, and sufficient amount of solid solution N cannot be obtained. On the other hand, when the cooling rate exceeds 300°C/sec, the material uniformity of the steel sheet in the width direction is insufficient. In addition, when the cooling stop temperature at a cooling rate of 10 to 300° C./sec after continuous annealing exceeds 500° C., refinement of the structure cannot be achieved.

Tempered rolling or leveling processing: elongation 0.5-10%

In the present invention, subsequent to the cold rolling and annealing process, temper rolling or leveling may be performed for the purpose of shape correction and smoothness adjustment. If the total elongation of temper rolling or leveling is less than 0.5%, the objects of shape correction and smoothness adjustment cannot be achieved. On the other hand, when it exceeds 10%, the ductility will fall. Furthermore, from the viewpoint of ensuring ductility, it is more preferably 5% or less. In addition, although temper rolling and leveling are different in processing form, it has been confirmed that there is no great difference in the effects of the two. Temper rolling and leveling are effective even after coating treatment.

The reasons for limiting the composition of the high-strength cold-rolled steel sheet of the seventh invention will be described.

C: 0.025 to 0.15%

C is an element that increases the strength of the steel sheet, and must be contained in an amount of 0.025% or more in order to uniformly and finely control the structure that is an important constitutional requirement of the present invention and ensure a sufficient amount of the martensite phase. On the other hand, when it exceeds 0.15%, the carbide fraction in the steel sheet is too large, and the ductility and formability are remarkably lowered. Also, as a more important problem, when the C content exceeds 0.15%, spot weldability, arc weldability, and the like are remarkably reduced. In this case, C is limited in the range of 0.025 to 0.15%. In addition, from the viewpoint of improving formability, it is preferably 0.08% or less. In addition, for applications requiring particularly good ductility, it is more preferably 0.05% or less.

Si: 1.0% or less

Si is a useful element that can increase the strength of the steel sheet without significantly lowering the ductility of the steel, and it is preferably contained in an amount of 0.005% or more, more preferably in an amount of 0.1% or more. On the other hand, Si greatly increases the transformation point during hot rolling, making it difficult to ensure the quality and shape, or also adversely affects the surface properties, chemical surface treatment, etc., especially the aesthetics of the steel sheet surface, and also affects the coating properties. Detrimental effects are limited to 1.0% or less in the present invention. If Si is 1.0% or less, the above-mentioned adverse effects can be suppressed to a low level. In addition, for applications requiring low levels of strength, especially surface aesthetics, Si is desirably 0.5% or less.

Mn: 2.0% or less

Mn is an effective element for preventing hot cracks caused by S, and it is better to add it according to the amount of S contained. In addition, Mn has a great effect on the refinement of crystal grains, and it is desirable to add it to improve the material. In addition, Mn is an extremely effective element for stably forming martensite during rapid cooling after continuous annealing. From the viewpoint of stably fixing S, it is desirable that Mn be contained in an amount of 0.2% or more. In addition, Mn is an element that increases the strength of the steel plate, and it is desirable to contain 1.2% or more when a super strength of TS 500MPa is required. More preferably 1.5% or more.

When the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet against fluctuations in hot rolling conditions, especially the dispersion of the strain age hardening characteristics are remarkably improved. However, when Mn is excessively contained over 2.0%, it is difficult to obtain a high r value which is one of the important requirements of the present invention, and the ductility is significantly lowered, so Mn is limited to 2.0% or less. For applications requiring better corrosion resistance and formability, the content is preferably 1.7% or less.

P: less than 0.08%

P is an element useful as a solid-solution strengthening element of steel, and it is preferably contained at 0.001% or more, more preferably 0.015% or more, from the viewpoint of increasing the strength. On the other hand, when it is contained excessively, the steel is embrittled, and the hemming workability of the steel sheet is also deteriorated. In addition, P has a strong tendency to segregate in steel, and thus causes embrittlement of the weld zone resulting therefrom. Therefore, P is limited to 0.08% or less. Furthermore, in applications where the hemming workability and weld toughness are particularly important, P is preferably 0.04% or less.

S: 0.02% or less

S is an element that exists in the form of inclusions in the steel sheet, reduces the ductility of the steel sheet, and also causes corrosion resistance to deteriorate. It is better to reduce it as much as possible. In the present invention, S is limited to 0.02% or less. Especially for applications requiring good workability, S is preferably 0.015% or less. In addition, when particularly excellent hemming workability is required, S is preferably 0.008% or less. Also, although the detailed mechanism is unclear, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening characteristics of the steel sheet at a high level.

Al: less than 0.02%

Al is an element that functions as a deoxidizer, improves the purity of steel, and refines the structure of the steel sheet. In the present invention, it is desirable to contain 0.001% or more. In the present invention, solid-solution N is used as a strengthening element, but an aluminum-killed steel containing an appropriate range of Al has better mechanical properties than a conventional ebullient steel that does not add Al. On the other hand, excessive Al content deteriorates the surface properties of the steel sheet and significantly reduces solid-solution N, making it difficult to obtain a large amount of strain age hardening. In this case, in the present invention, Al is limited to 0.02% or less. Furthermore, from the viewpoint of material stability, Al is preferably 0.001 to 0.015%. In addition, reducing the Al content may also cause grain coarsening, but in the present invention, this situation is effectively prevented by limiting other alloy elements to the optimum amount and making the annealing conditions within the optimum range.

N: 0.0050～0.0250%

N is an element that increases the strength of a steel sheet through solid solution strengthening and strain age hardening, and is the most important element in the present invention. In addition, in the present invention, by containing an appropriate amount of N, adjusting the Al content to an appropriate value as described above, and controlling manufacturing conditions such as hot rolling conditions and annealing conditions, the necessary and stable N in cold-rolled products or coated products can be ensured. Sufficient N in solid solution. Accordingly, the effect of increasing the strength (yield stress and tensile strength) by solid solution strengthening and strain age hardening can be fully exerted, and the following target values of the mechanical properties of the steel sheet of the present invention can be stably obtained: the tensile strength is 440 MPa or more, The bake hardening amount (BH amount) is 80MPa or more, and the tensile strength increase ΔTS before and after strain aging treatment is 40MPa or more.

When N is less than 0.0050%, it is difficult to stably exhibit the aforementioned effect of increasing the strength. On the other hand, when N exceeds 0.0250%, the occurrence rate of internal defects in the steel sheet increases, and slab cracks during continuous casting often occur. Therefore, N is limited to the range of 0.0050 to 0.0250%. Furthermore, N is preferably 0.0070 to 0.0170% from the viewpoint of improving the stability of the material and the yield in consideration of the overall manufacturing process. In addition, as long as the amount of N is within the range of the present invention, there is no adverse effect on weldability or the like at all.

Solid solution N: 0.0010% or more

In cold-rolled products, in order to ensure sufficient strength and effectively exert strain age hardening caused by N, at least 0.0010% or more of solid-solution N (also called solid-solution N) must exist in the steel sheet.

Here, the amount of solid solution N is the amount of precipitated N subtracted from the total amount of N in the steel, and the obtained value is taken as the amount of solid solution N. Furthermore, the present inventors compared and studied various methods as an analysis method for the amount of precipitated N, and as a result, it is effective to obtain it by an electrolytic extraction method using a constant potential electrolysis method. In addition, as a method of dissolving matrix iron for extraction analysis, there are acid decomposition method, halogen method, and electrolysis method. Among them, the electrolytic method does not decompose extremely unstable precipitates such as carbides and nitrides, and can only dissolve matrix iron stably. An acetylacetone system was used as the electrolytic solution, and it was electrolyzed at a constant potential. In the present invention, the results of measuring the amount of precipitated N using the constant potential electrolysis method showed a good correspondence with actual material changes.

In view of this, in the present invention, the residue extracted by the constant potential electrolysis method is chemically analyzed to determine the amount of N in the residue, which is taken as the amount of precipitated N.

Furthermore, when a higher amount of BH and ΔTS are required, it is preferable to make the amount of solid solution N 0.0020% or more, and to obtain a higher value, it is preferable to make it 0.0030% or more. The upper limit of the amount of solid solution N is not particularly limited, but even if the total amount of N added remains as a whole, the decrease in mechanical properties is small.

N/Al (ratio of N content to Al content): 0.3 or more

In order to stably leave 0.0010% or more of solid-solution N in the product state, the amount of Al, which is an element that strongly fixes N, must be limited. As a result of examining steel sheets in which combinations of N content (0.0050 to 0.0250%) and Al content (0.02% or less) were widely varied within the composition range of the present invention, it was found that by setting N/Al to 0.3 or more, The solid solution N in cold-rolled products and coated products is stably maintained at 0.0010% or more. Therefore, N/Al is limited to 0.3 or more.

In the present invention, in addition to the above-mentioned composition, it is preferable to further contain one or more of the following groups d to g, wherein,

Group d: 1 or more of Cu, Ni, Cr, and Mo, with a total amount of 1.0% or less;

Group e: one or more of Nb, Ti, V, the total amount is less than 0.1%;

Group f: B below 0.0030%;

Group g: 1 or 2 kinds of Ca and REM, the total amount is 0.0010-0.010%.

Group d elements: Cu, Ni, Cr, and Mo are all elements that contribute to the increase in the strength of the steel sheet, and can be selected and contained alone or in combination as needed. This effect is seen when Cu, Ni, Cr, and Mo are contained in an amount of 0.005% or more. However, when the content is too large, the thermal deformation resistance increases, or the chemical surface treatability and generalized surface treatment characteristics deteriorate, and the weld zone hardens, and the weld zone formability deteriorates. In addition, the r value also tends to decrease. Therefore, the total amount of elements in group a is preferably 1.0% or less. Furthermore, when Mo is contained in a large amount of 0.05% or more, the r value may be significantly lowered. When Mo is contained in the present invention, it is preferably limited to less than 0.05%.

Group e elements: Nb, Ti, and V are all elements that contribute to grain refinement and homogenization, and can be selected and contained singly or in combination as needed. This effect is seen when each of Nb, Ti, and V is contained at 0.005% or more. However, when the content is too large, the thermal deformation resistance increases while chemical surface treatment and surface treatment properties in a broad sense deteriorate. Therefore, the total amount of group b elements is preferably 0.1% or less.

Group f elements: B is an element having an effect of improving the hardenability of steel, and may be contained as necessary for the purpose of increasing the fraction of low-temperature transformation phases other than the ferrite phase and increasing the strength of the steel. This effect is seen when B is contained in an amount of 0.0005% or more. However, if the equivalent is too large, the thermal deformation ability decreases, and the solid solution N decreases due to the formation of BN. Therefore, B is preferably 0.0030% or less.

Group g elements: Ca and REM are elements that play a role in controlling the shape of inclusions, especially when curling formability is required, it is better to contain them alone or in combination. At this time, when the total amount of d group elements is less than 0.0010%, the effect of controlling the morphology of inclusions is insufficient, while on the other hand, when it exceeds 0.010%, the occurrence of surface defects becomes significant. Therefore, it is preferable to limit the total amount of the d group elements to the range of 0.0010 to 0.010%. According to this, the hemming workability can be improved without being accompanied by the occurrence of surface defects.

Next, the structure of the steel sheet of the present invention will be described.

Area ratio of ferrite phase: 80% or more

The cold-rolled steel sheet of the present invention is aimed at steel sheets for automobiles that require a certain degree of workability. In order to ensure ductility, a structure containing 80% or more of the ferrite phase in terms of area ratio is formed. When the area of the ferrite phase When the ratio is less than 80%, it is difficult to ensure the ductility required for a steel sheet for automobiles requiring workability. Furthermore, when better ductility is required, the area ratio of the ferrite phase is desirably 85% or more. In addition, the ferrite referred to in the present invention refers to so-called polygonal ferrite that does not remain in a strained state.

Average grain size of ferrite phase: 10 μm or less

In the present invention, as the average crystal grain size, the value calculated from the cross-sectional structure photograph by the quadrature method prescribed by ASTM and the nominal diameter obtained by the cutting method prescribed by the same ASTM (for example, refer to Umemoto et al.: heat treatment , 24(1984), 334), whichever is larger.

The cold-rolled steel sheet of the present invention ensures a predetermined amount of solid-solution N in the product stage, but according to the experiments and studies of the present inventors, even in steel sheets having the same amount of solid-solution N, the strain-age hardening characteristics sometimes vary. , it has been found that one of the main factors is the grain size. In the structure of the present invention, by making the average grain size at least 10 μm or less, preferably 8 μm or less, a high BH amount and ΔTS can be stably obtained. Although the detailed mechanism is unclear, it can be speculated that it is related to the segregation and precipitation of alloy elements to the grain boundaries, and the influence of processing and heat history on them.

Therefore, in order to stabilize the strain age hardening characteristics, the average crystal grain size of the ferrite phase must be 10 μm or less, preferably 8 μm or less.

As described above, in order to ensure the ductility as a steel sheet for automobiles and to stabilize the strain age hardening characteristics, in the present invention, ferrite containing 80% or more of the average grain size in terms of area ratio is formed. organization.

Area ratio of martensite phase: 2% or more

The cold-rolled steel sheet of the present invention contains 2% or more of martensite as the second phase in terms of area ratio. By the presence of 2% or more of martensite, good ductility and greater strain age hardenability can be obtained. Although the detailed mechanism is not clear, it is presumed that due to the presence of the martensitic phase, strain is stored inside very effectively during pre-strain processing before aging. In addition, the presence of the martensite phase is also effective in improving aging deterioration. Furthermore, in order to obtain a good strength-ductility balance and a low yield ratio, the area ratio of the martensite phase is preferably 5% or more. Furthermore, the presence of a large amount of martensite phase exceeding 20% has a problem of lowering ductility. Therefore, the area ratio of the martensite phase is 2% or more, preferably 5% or more and 20% or less.

As the second phase, in addition to the above-mentioned martensite phase, there is no problem with pearlite, bainite, and retained austenite. However, in the present invention, the ferrite phase fraction must be 80% or more and the martensite Tensile phase fraction is 2% or more. The total area ratio of pearlite, bainite, and retained austenite is limited to less than 18%.

The cold-rolled steel sheet of the present invention having the above-mentioned composition and structure has a tensile strength (TS) of 440 MPa or more and about 780 MPa or less, and has a high r value of 1.2 or more by controlling the aggregate structure of parent phase ferrite, Cold rolled steel sheet with excellent strain age hardening properties. Steel plates with TS less than 440MPa cannot be widely used in components with structural component elements. In addition, in order to further expand the applicable range, it is desirable that TS is 500 MPa or more. In addition, if the r value is less than 1.2, it cannot be applied to a wide range of press-formed parts. In addition, the preferable range of r value is 1.4 or more.

In the present invention, the so-called "excellent strain age hardening properties" means: as mentioned above, after pre-deformation with a tensile strain of 5%, when aging treatment is performed at a temperature of 170°C and kept for 20 minutes, the aging The deformation stress increase before and after treatment (denoted as BH amount; BH amount=yield stress after aging treatment-pre-deformation stress before aging treatment) is more than 80MPa, and the amount before and after strain aging treatment (predeformation+aforementioned aging treatment) The increase in tensile strength (denoted as ΔTS; ΔTS=tensile strength after aging treatment-tensile strength before pre-deformation) is above 40 MPa.

When specifying the strain age hardening characteristics, the amount of prestrain (predeformation) becomes an important factor. The inventors of the present invention have studied the influence of the pre-strain amount on the strain age hardening characteristics assuming that the deformation pattern is applicable to steel sheets for automobiles. It can be sorted by the amount of strain (tensile strain) equivalent to 1 axis; ②In actual components, the amount of strain equivalent to 1 axis is about more than 5%; ③The relationship between the strength of the component and the strain aging at 5% of the pre-strain has been found out The intensities (YS and TS) obtained after processing correspond well. Based on this knowledge, the present invention sets the pre-deformation of the strain aging treatment as a tensile strain of 5%.

Conventional painting and baking conditions have been adopted as a standard of 170°C x 20 minutes. Furthermore, when a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solid solution N, hardening can be achieved even with a milder (low temperature side) treatment, in other words, aging conditions can be obtained in a wider range. . In addition, in general, in order to strive for the amount of hardening, it is advantageous to keep at a higher temperature for a longer time within the limit of not softening due to excessive aging.

Specifically, the lower limit of the heating temperature at which hardening becomes significant after pre-deformation of the steel sheet of the present invention is about 100°C. On the other hand, when the heating temperature exceeds 300°C, the hardening reaches its peak, and on the contrary, there is a tendency to slightly soften, and furthermore, the generation of thermal strain and temper color becomes remarkable. In addition, regarding the holding time, when the heating temperature is about 200° C., it is about 30 seconds or more, and sufficient curing can basically be achieved. In order to obtain greater stable hardening, the holding time is preferably 60 seconds or more. However, when the holding time exceeds 20 minutes, not only further hardening cannot be obtained, but also the production efficiency is significantly lowered, which is disadvantageous in practical use.

Therefore, in the present invention, as the aging treatment conditions, the heating temperature of 170° C., which is the conventional paint baking treatment condition, was evaluated for 20 minutes, and the holding time was set. Even under the aging treatment conditions of low-temperature heating and short-time holding that the conventional paint-bake type steel sheet cannot achieve sufficient hardening, the steel sheet of the present invention can stably achieve large hardening. Furthermore, the method of heating is not particularly limited, and in addition to the atmosphere heating carried out by a furnace used in general paint baking, for example, heating by induction heating, non-oxidizing flame, laser, plasma, etc. can be carried out very well. Use it well.

The strength of automotive parts must be able to resist complex stress loads from the outside, so the strength characteristics of the raw steel sheet are not only important in the small strain region, but also in the large strain region. Based on this point, the present inventors set the BH amount of 80 MPa or more and the ΔTS amount of 40 MPa or more in the steel sheet of the present invention to be used as a material for automobile parts. Furthermore, it is more preferable that the amount of BH is 100 MPa or more, and ΔTS is 50 MPa or more. In addition, by setting the heating temperature during the aging treatment to a higher temperature and/or setting the holding time to a longer time, the amount of BH and the amount of ΔTS can be further increased.

In addition, the steel sheet of the present invention has the advantage that a strength increase of about 40% of that of complete aging can be expected only after being left at room temperature for about one week without special heating after forming.

In addition, the steel sheet of the present invention also has the advantages provided by conventional aging steel sheets: in the unformed state, aging deterioration does not occur even if left at room temperature for a long time (YS increases and E1 (elongation) decreases Phenomenon). Furthermore, in order not to cause problems in actual stamping, the 3-month aging at room temperature before stamping must reach: the increase in YS is 30 MPa or less, the decrease in elongation is 2% or less, and the yield The recovery of point elongation was 0.2% or less.

In addition, in the present invention, there is no problem even if electroplating or hot-dipping is performed on the surface of the above-mentioned cold-rolled steel sheet of the present invention. These coated steel sheets also showed the same levels of TS, BH, and ΔTS as those before plating. As the type of coating, electro-galvanizing, hot-dip galvanizing, alloyed hot-dip galvanizing, electro-tin plating, electro-chrome plating, electro-nickel plating, etc. can be favorably used.

The manufacturing method of the steel plate of the 8th this invention is demonstrated.

The steel plate of the present invention is basically manufactured by sequentially performing the following steps: heating a steel slab having the composition in the above-mentioned range, rough rolling it to form a thin slab, subjecting the thin slab to finish rolling, cooling after finishing rolling, and coiling to form a thin slab. Hot-rolling process of hot-rolled sheet, cold-rolling process of pickling and cold-rolling hot-rolled sheet to produce cold-rolled sheet, cold-rolled sheet annealing process of subjecting cold-rolled sheet to box annealing, followed by continuous annealing .

The slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macrosegregation of components, but it may also be produced by an ingot casting method or a thin slab casting method. In addition, after manufacturing the slab, in addition to the conventional method of once cooling to room temperature and then reheating, straight rolling is carried out in which the hot slab is placed in a heating furnace without cooling and rolled, or is directly heated after being slightly heated. Energy-saving processes such as rolling direct rolling can also be used without problems. In particular, in-line rolling is one of useful techniques for effectively securing N in a solid solution state.

First, the reasons for limiting the conditions of the hot rolling process will be described.

Slab heating temperature: above 1000°C

In the initial state of hot rolling, in order to ensure a necessary and sufficient amount of solid solution N and to make the amount of solid solution N in the product meet the target value, the heating temperature of the slab is preferably set at 1000° C. or higher. Furthermore, since the loss increases as the oxidation weight increases, it is desirable to set it at 1280°C or lower.

The slab heated under the above conditions is rough rolled into a thin slab. In addition, the conditions of rough rolling do not need to be specifically defined, What is necessary is just to carry out by a conventional method. However, from the viewpoint of securing the amount of solid-solution N, it is desirable to perform the process in as short a time as possible. The thin slab is then finish-rolled to make hot-rolled plate.

Furthermore, in the present invention, between rough rolling and finish rolling, it is desirable to join front and rear adjacent thin slabs together and roll them continuously. It is preferable to use a pressure welding method, a laser welding method, an electron beam welding method, etc. as a joining means.

Continuous rolling eliminates so-called unsteady portions of rolling at the front and rear ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length and width of the coil (material to be processed). This is extremely effective in improving not only the hot-rolled steel sheet but also the cross-sectional shape and size of the cold-rolled steel sheet. In addition, tension can always be applied when cooling on the hot metal roller table after rolling, so the shape of the steel sheet can be maintained well.

In addition, in order to stably pass the front end of the coil by performing continuous rolling, lubricated rolling, which cannot be used due to the problems of plate passability and biting property, can be used in the usual single rolling of each thin slab. Accordingly, the rolling load can be reduced, and at the same time, the surface pressure of the roll can be reduced, and the life of the roll can be extended.

In addition, in the present invention, on the feed side of the finish rolling mill between rough rolling and finish rolling, a sheet slab edge heater for heating the transverse end of the slab, a sheet slab heater for heating the longitudinal end of the slab are used. Either one or both sides, it is better to make the temperature distribution in the width direction and longitudinal direction of the thin slab uniform. According to this, material dispersion in the steel sheet can be further reduced. The thin slab edge heater and the thin slab heater are preferably induction heating.

The order of use is expected to firstly use thin slab edge heaters to compensate for temperature differences in the width direction. The amount of heating at this time is also based on the steel composition, but it is better to set the temperature distribution range in the width direction on the side of the finish rolling to about 20°C or less. The temperature difference in the length direction is then compensated by the thin slab heater. The amount of heating at this time is set so that the temperature at the longitudinal end is about 20°C higher than the temperature in the central region.

Finishing rolling temperature: above 800°C

In order to obtain a uniform and fine hot-rolled mother plate structure, the finishing rolling temperature FDT is set to be above 800°C. When the FDT is lower than 800°C, the structure of the steel plate becomes inhomogeneous, and the processed structure remains partially, and the inhomogeneity of the structure cannot be eliminated after the cold rolling and annealing process, and remains. Therefore, when a high winding temperature is used to avoid the residue of the processed structure, coarse crystal grains are generated, and the same problem occurs. In addition, when the coiling temperature is high, the amount of solid-solution N significantly decreases, so it is difficult to obtain the target tensile strength of 440 MPa or more. Due to this situation, the finishing rolling temperature FDT is taken as 800°C. In order to further improve the mechanical properties, it is desirable to set the FDT to 820° C. or higher. The upper limit of FDT is not particularly defined, but when it is excessively high, occurrence of scale defects and the like becomes remarkable. Furthermore, FDT is preferably set to about 1000°C.

In addition, cooling after finish rolling is not particularly strictly limited, but it is desirable to satisfy the following conditions in terms of material uniformity in the longitudinal and transverse directions of the steel sheet. That is, in the present invention, it is desirable to start cooling immediately (within 0.5 seconds) after finishing rolling, and to set the average cooling rate during cooling to 40° C./second or more. By satisfying this condition, the high-temperature region where AlN is precipitated can be rapidly cooled, and N in a solid solution state can be effectively secured. When the cooling start time or the cooling rate does not satisfy the above conditions, the grain growth proceeds excessively, making it difficult to achieve grain refinement, and there is a tendency for AlN precipitation due to the strain energy introduced by rolling to be promoted, which may Lack of solid solution N, there is a tendency for the structure to be uneven. Furthermore, from the viewpoint of ensuring the uniformity of material and shape, it is preferable to suppress the cooling rate to 300°C/sec or less.

Winding temperature: below 800°C

As the coiling temperature CT decreases, the strength of the steel sheet tends to increase. In order to ensure the target tensile strength TS440MPa or more, it is better to take CT below 800°C. Furthermore, when the CT is less than 200° C., the shape of the steel sheet tends to be uneven, and there is a high risk of occurrence of defects in actual operation, showing a tendency for the uniformity of the material to decrease. Therefore, it is desirable that CT be set at 200°C or higher. Furthermore, in the case where material uniformity is more required, it is better to take CT above 300°C. More preferably, it is 350° C. or higher. In addition, in the present invention, lubricated rolling may be performed in order to reduce the hot rolling load during finish rolling. There is an effect of making the shape and material of the hot-rolled sheet more uniform by performing lubricated rolling. Furthermore, the friction coefficient during lubricated rolling is preferably in the range of 0.25 to 0.10. In addition, the operation of hot rolling is stabilized by combining lubricated rolling and continuous rolling.

The hot-rolled sheet subjected to the above-mentioned hot-rolling step is then subjected to pickling and cold-rolling in the cold-rolling step to be a cold-rolled sheet.

The pickling conditions may be generally known conditions and are not particularly limited. Furthermore, when the scales of the hot-rolled sheet are extremely thin, cold rolling may be performed without pickling.

In addition, the cooling conditions may be generally known conditions, and are not particularly limited. Furthermore, from the viewpoint of ensuring the uniformity of the structure, the cold reduction rate is preferably 40% or more. The reasons for limiting the conditions of the cold rolling process will be described below.

The cold-rolled sheet is then subjected to a cold-rolled sheet annealing process consisting of box annealing and continuous annealing.

Box annealing temperature: above recrystallization temperature to below 800°C

In the present invention, the cold-rolled sheet is subjected to box annealing to control the microstructure of the ferrite phase that becomes the matrix. By controlling the microstructure of the ferrite phase, a high r-value of the product plate can be achieved. This box annealing facilitates the formation of the (111) aggregate structure desired for realizing a high r value in the product sheet.

When the box annealing temperature is lower than the recrystallization temperature, the recrystallization cannot be completed, the microstructure of the ferrite phase cannot be adjusted, and a high r value cannot be obtained. On the other hand, when box annealing is performed at a temperature exceeding 800° C., the occurrence of surface defects on the steel sheet becomes remarkable, and the initial purpose cannot be achieved. It should be noted that the box annealing is preferably carried out in an annealing atmosphere containing mainly nitrogen gas and containing 3 to 5% hydrogen gas. In this case, the heating and cooling rate may be about 30° C./hour under normal box annealing conditions. In addition, by making the annealing atmosphere gas 100% hydrogen, faster heating and cooling rates can be employed.

Continuous annealing temperature: above Ac ₁ transformation point to (Ac ₃ transformation point -20°C) below

When the continuous annealing temperature is lower than the Ac ₁ transformation point, the martensite phase cannot be formed after annealing. On the other hand, when it exceeds (Ac ₃ transformation point-20℃), the desired assembly formed in the box annealing The structure is lost by the phase transformation, so a finished sheet with a high r value cannot be obtained. Therefore, the continuous annealing temperature is preferably above the Ac ₁ transformation point and below (Ac ₃ transformation point-20°C). In addition, the retention time of the continuous annealing time is preferably as short as possible from the viewpoint of production efficiency, microstructure refinement, and securing the amount of solid solution N. On the other hand, from the viewpoint of the stability of the work, the holding time is preferably 10 seconds or more. In addition, from the viewpoint of refining the structure and ensuring the amount of solid solution N, it is preferable to set it at 120 seconds or less. Furthermore, from the viewpoint of stabilization of the material, it is more preferably 20 seconds or more.

Cooling after continuous annealing: cooling to a temperature range below 500°C at a cooling rate of 10 to 300°C/sec

Cooling after soaking in continuous annealing is important from the viewpoint of refining the structure, forming a martensitic phase, and securing the amount of solid solution N. In the present invention, cooling is performed continuously to a temperature range below 500° C. at a cooling rate of at least 10° C./sec. When the cooling rate is less than 10°C/sec, the necessary amount of martensite, uniform and fine structure, and sufficient amount of solid solution N cannot be obtained. On the other hand, when the cooling rate exceeds 300°C/sec, the amount of supersaturated solid-solution C increases remarkably, so that the material uniformity of the steel sheet in the transverse direction decreases. When the cooling stop temperature at a cooling rate of 10 to 300°C/sec after continuous annealing exceeds 500°C, refinement of the structure cannot be achieved.

Conditions for overaging treatment: Continuing with cooling after continuous annealing, the residence time in the temperature zone below the cooling stop temperature of the cooling and above 350°C is 20 seconds or more

Successively to the cooling stop after soaking in the continuous annealing, the overaging treatment with a residence time of 20 seconds or more may be performed in a temperature range between the cooling stop temperature and 350° C. or higher. By performing the overaging treatment, the amount of solid solution C can be selectively reduced while maintaining the amount of solid solution N. If the retention temperature range is less than 350°C, it will take a long time to reduce the solid solution C, resulting in a decrease in productivity, so it is preferable to set the temperature range to 350°C or higher.

By staying in a temperature range of not lower than the cooling stop temperature but not lower than 350° C. for 20 seconds or more, the amount of solid solution C can be reduced, and a higher degree of non-aging at room temperature can be achieved. Further improvement can be expected by making the residence time longer, but the effect tends to be saturated at about 120 seconds, so the residence time is preferably 120 seconds or less.

In order to obtain a large amount of strain age hardening, it is advantageous to use both solid solution C and solid solution N, but when solid solution C is used, aging deterioration at room temperature becomes significant, and the applicable parts of the steel sheet are limited. Therefore, in order to manufacture a strain-age-hardening type steel sheet with wide applicability, it is preferable to perform an overaging treatment after securing a sufficient amount of solid-solution N.

Furthermore, in the case of manufacturing a high-strength cold-rolled steel sheet having a hot-dip coating on the surface of the high-strength cold-rolled steel sheet of the present invention, continuous annealing is carried out immediately after box annealing on the continuous hot-dip coating production line, and cooling after continuous annealing Subsequently, hot-dip galvanizing or alloying treatment can be performed to manufacture hot-dip galvanized steel sheets.

Tempered rolling or leveling processing: elongation 0.2 to 15%

In the present invention, subsequent to the cold rolling and annealing process, temper rolling or leveling may be performed for the purpose of shape correction and smoothness adjustment. If the total amount of elongation in temper rolling or leveling is less than 0.2%, the intended purpose of shape correction and smoothness adjustment cannot be achieved. On the other hand, when it exceeds 15%, it causes a remarkable decrease in ductility. Furthermore, although temper rolling and leveling are different in processing form, it has been confirmed that there is no great difference in the effects of both. Temper rolling and leveling are effective even after coating treatment.

Hereinafter, for reference, forming conditions when the steel sheet of the present invention is used for forming such as press forming and subsequent heat treatment conditions for increasing strength will be described. When the steel sheet of the present invention is used for pressing such as drawing, for example, the strain introduced by the pressing is several percent to several tens of percent. The amount of strain varies depending on the formed part, but about 5 to 10% of strain can be introduced into inner panels and structural parts in the automotive field.

Next, these formed parts are subjected to heat treatment such as painting and baking treatment, but the steel plate of this invention can effectively increase the strength of the formed product after heat treatment. Furthermore, in this invention, as a method of evaluating the bake hardenability in the laboratory, a tensile test piece of JIS No. 5 size is prepared along the rolling direction, and a tensile strain of 10% is given by a tensile testing machine. , and then after applying the heat treatment, the tensile test was performed again. In particular, when evaluating the properties after heat treatment in the low temperature range, the heat treatment conditions were set at 120° C. for 20 minutes. This test evaluates the properties of a finished part that has been heat treated immediately after stamping.

That is, in this invention, the difference (ΔTS) between the tensile strength after such a tensile strain-imparting heat treatment and the tensile strength of the product is defined as the strength-enhancing heat-treatability.

Generally, in order to increase the strength increase of a molded product, it is better to have a large amount of strain introduced by forming or a high heat treatment temperature after processing.

However, when the amount of strain imparted to the steel sheet of this invention is about 5 to 10% as described above, a sufficient increase in strength can be achieved even if the heat treatment temperature after forming is lower than conventional ones, that is, the heat treatment temperature is 200°C or lower. However, when the heat treatment temperature is lower than 120° C., a sufficient effect of increasing the strength cannot be obtained when the strain is low. On the other hand, when the heat treatment temperature after forming is a temperature higher than 350° C., softening will proceed. Therefore, the heat treatment temperature after forming is preferably about 120 to 350°C.

In addition, as the heating method, methods such as hot air heating, infrared furnace heating, hot bath heat treatment, energization heating, and high-frequency heating can be used, and are not particularly limited. Alternatively, only the portion where the strength is to be increased may be selectively heated.

Example

In the following examples, the amount of solid solution N, microstructure, tensile properties, r value measurement, strain age hardening properties, and aging properties were investigated. The survey method is as follows.

(1) Solid solution N content

The amount of solid solution N is obtained by subtracting the amount of precipitated N from the total amount of N in the steel obtained by chemical analysis. Here, the amount of precipitated N was obtained by the analysis method using the above-mentioned constant potential electrolysis method.

(2) Microstructure

Take a test piece from each cold-rolled and annealed sheet, use an optical microscope or a scanning electron microscope to photograph the microstructure of a cross-section (C-section) perpendicular to the rolling direction, and use an image analysis device to obtain the structure fraction of ferrite and the type and tissue fraction of phase 2.

(3) Grain size

In the present invention, as the crystal grain size, the value obtained from the cross-sectional structure photograph by the quadrature method specified by ASTM and the nominal diameter obtained from the cross-sectional structure photograph by the cutting method specified by ASTM (for example, refer to Umemoto etc.: heat treatment, 24(1984), 334), a certain large value.

(4) Tensile properties

JIS No. 5 test pieces are prepared from each cold-rolled and annealed sheet along the rolling direction, and tensile tests are performed at a strain rate of 3×10 ^-3 /sec in accordance with JIS Z 2241 to obtain yield stress YS and tensile strength TS , Elongation E1.

(5) Strain age hardening characteristics

A JIS No. 5 test piece is prepared from each cold-rolled annealed sheet along the rolling direction, and a tensile prestrain of 5% is applied here as a pre-deformation, followed by a heat treatment equivalent to a paint-bake treatment at 170°C for 20 minutes. , implement a tensile test at a strain rate of 3×10 ^-3 /sec, obtain the tensile properties (yield stress YSBH, tensile strength TSBH) after the pre-deformation-painting and baking treatment, and calculate the BH amount=YSBH-YS5%, ΔTS=TSBH-TS. Furthermore, YS5% is the deformation stress when the product plate is pre-deformed by 5%, YSBH and TSBH are the yield stress and tensile strength after pre-deformation-painting and baking treatment, and TS is the tensile strength of the product plate.

(6) Determination of r value

JIS No. 5 test pieces were prepared from the rolling direction (L direction), the direction 45° to the rolling direction (D direction), and the direction 90° to the rolling direction (C direction) of each cold-rolled and annealed sheet. The transverse strain and the plate thickness direction strain of each test piece when a 5% uniaxial tensile strain is applied to these test pieces are obtained, and the ratio of the transverse strain and the plate thickness direction strain as the definition formula of the r value is obtained.

r=ln(w/w ₀ )/In(t/t ₀ )

(where, w ₀ and t ₀ are the width and thickness of the test piece before the test, and w and t are the width and thickness of the test piece after the test.) Calculate the r value in each direction, and use the following formula

_raverage =(r _L +2r _D +r _C )/4 Calculate the average r value raverage. Among them, r _L is the r value in the rolling direction (L direction), r _D is the r value in the direction (D direction) at 45° to the rolling direction (L direction), and r _C is the r value at 90° to the rolling direction The r value of the direction (C direction). In addition, in order to improve the accuracy of the experiment, it is assumed that the volume is constant, and it is calculated from changes in tensile strain and lateral strain.

(7) Aging characteristics

A JIS No. 5 test piece was prepared from each cold-rolled and annealed sheet, and the test piece was subjected to an aging treatment at 50°C for 200 hours, and then a tensile test was performed. From the obtained results, the difference in yield elongation before and after the aging treatment ΔY- E1, evaluating the aging characteristics at room temperature. When ΔY-E1 is 0, it is evaluated as non-aging property that it is excellent in room temperature aging resistance.

(8) Forming - tensile strength after heat treatment

The tensile strength after forming-heat treatment is obtained from the product plate along the in-process direction of the JIS No. 5 test piece, and after giving a 10% pre-strain, it is 120°C and 170°C, which is the heat treatment temperature equivalent to the paint baking in the past. It was obtained by performing heat treatment at °C for 20 minutes and measuring the tensile strength.

(9) Decrease in total elongation due to aging at room temperature (ΔE1)

The decrease in total elongation due to aging at room temperature (ΔE1) is the total elongation measured by taking a JIS No. 5 test piece from the product plate along the rolling direction and using another JIS 5 test piece taken along the rolling direction. It was obtained from the difference in the total elongation measured after the accelerated treatment of room temperature aging (100° C., 8 hours) was applied to the No. test piece.

Example 1

The steel slab with the composition shown in Table 1 is made into a hot-rolled sheet with a thickness of 3.5mm under the conditions shown in Table 2, and then made into a cold-rolled sheet with a thickness of 0.7mm. —Recrystallization annealing is performed on the alloying hot-dip galvanizing line, and then the alloying hot-dip galvanizing treatment is performed, and then the reduction ratio: 1.0% quenched and tempered rolling is performed to manufacture cold-rolled steel sheets and the average coating amount on one side : 45g/m2 Alloyed hot-dip galvanized steel sheet coated on both sides. In addition, No. 3 and 8 in Table 2 had hot rolling completion temperatures below the Ar3 transformation point, and other than that, they were above the Ar3 transformation point.

Table 3 shows the results of investigations on the tensile strength and r-value of the cold-rolled steel sheets and galvannealed steel sheets thus obtained, and changes in tensile strength after forming and heat treatment.

It can be seen from Table 3 that the cold-rolled steel sheet and alloyed hot-dip galvanized steel sheet obtained according to the invention can obtain high r value and excellent strain age hardening property compared with the comparative example. In addition, especially among the suitable examples, the amount of reduction in elongation due to aging at room temperature in the steel sheet having a grain size of 20 μm or less was also reduced to 2.0% or less as expressed by ΔE1.

Example 2

Using the slab of steel No. B described in Table 1, after hot rolling at the slab heating temperature: 1100°C and hot finish rolling temperature: 900°C under the same manufacturing conditions as No. 2 in Table 2, the coiling temperature : Coiled into coils at 550°C. After the coil was cold-rolled at a rolling reduction of 80%, recrystallization annealing was performed at 840°C. The product properties of the obtained cold-rolled steel sheet were tensile strength TS = 365 MPa, r value = 1.7. Take J1S No. 5 test pieces along the rolling direction from this cold-rolled plate, after utilizing the tensile testing machine to give 10% tensile strain, implement heat treatment under the heat treatment conditions (temperature, time) shown in Table 4, again Do a tensile test. Table 4 also records the amount of increase in tensile strength (ΔTS) from the tensile strength (TS=365 MPa) of the product before straining.

As shown in Table 4, the higher the heat treatment temperature and the longer the heat treatment time, the greater the increase in strength. However, even if the heat treatment temperature of the inventive steel plate is as low as 120°C and the holding time is as short as 2 minutes, sufficient strength of 82 MPa can be obtained. The amount of increase in tensile strength (more than 85% during 20 minutes of heat treatment) can obtain good strain age hardening properties even with low temperature and short time heat treatment. Furthermore, in structural members of automobiles, etc., there is no problem in performing heat treatment at normal temperature and time in order to obtain a stable strength-enhancing effect. It was also shown that the same results as those in Table 4 were obtained for the obtained galvannealed steel sheet by hot-dip galvanizing and heat alloying treatment on the cold-rolled steel sheet.

Example 3

Steel slabs having the compositions shown in Table 6 were hot-rolled under the conditions shown in Table 7 to form hot-rolled sheets with a thickness of 3.5 mm. These hot-rolled sheets were cold-rolled under the conditions shown in Table 7 to produce cold-rolled sheets with a thickness of 0.7 mm. These cold-rolled sheets were recrystallized and annealed under the conditions shown in the table, and some of them were recrystallized and annealed as shown in the table. Hot-dip galvanizing or alloyed hot-dip galvanizing under certain conditions. The solid solution N content, microstructure, tensile properties, and strain age hardening properties of the obtained product sheets were investigated.

The results are shown in Table 8. It can be seen from the table that all the steel sheets manufactured according to the present invention satisfy TS×value ≥ 750MPa (when B is added in combination with one or more of Nb, Ti, V, further TS×r value ≥ 850MPa), BH≧80 MPa and ΔTS≧40 MPa, however, in the comparative example, one or more of these three properties did not reach the level of the present invention.

Example 4

Embodiments of the present invention will be described below.

The molten steel with the composition shown in Table 9 was smelted in a converter, and a billet was made by continuous casting. These steel billets were heated under the conditions shown in Table 10, rough rolled into thin slabs, and then subjected to a hot rolling process of finish rolling under the conditions shown in Table 10 to make hot rolled sheets. In addition, the Ar3 transformation point was measured using the processing transformation measuring apparatus (manufactured by Fuji Denpa Industrial Machinery Co., Ltd.) under simulated conditions, and Table 10 shows the hot finish rolling conditions.

These hot-rolled sheets were made into cold-rolled sheets through a cold-rolling process consisting of pickling and cold-rolling under the conditions shown in Table 10. Next, these cold-rolled sheets were subjected to continuous annealing under the conditions shown in Table 10. Some of them are subjected to temper rolling after the cold rolling and annealing process.

Regarding the obtained cold-rolled annealed sheet, the solid solution N content, microstructure, tensile properties, r-value measurement, strain age hardening properties, and aging properties were investigated.

Furthermore, hot-dip galvanizing was performed on the surface of the steel sheets of No. 4 and No. 10 to form plated steel sheets, and various characteristics were evaluated in the same manner.

These results are shown in Table 11.

In the examples of the present invention, all exhibit excellent ductility, particularly high BH content and ΔTS, have excellent strain age hardening properties, have a high r value with an average r value of 1.2 or more, and are excellent in non-aging properties during aging at room temperature. Resistance to aging at room temperature. Furthermore, the properties of coated steel sheets obtained by hot-dip galvanizing on the surface of No. 4 and No. 10 steel sheets are restrained by the transverse shrinkage of the coating layer, so compared with cold-rolled steel sheets, the average r value is reduced by 0.2, and the elongation E1 is reduced. About 1%, but the strain age hardenability, room temperature aging resistance and the characteristic ratio before plating basically do not change. On the contrary, the comparative examples outside the scope of the present invention have poor ductility, little BH amount, ΔTS, or significant aging deterioration, and do not have all the target properties, and cannot be said to be steel sheets with sufficient properties.

C, Al, N, and N/Al of steel sheet No. 11 deviated from the range of the present invention, so the r value, BH amount, ΔTS, and room temperature aging resistance decreased. In addition, B and Nb of steel plate No. 12 deviate from the range of the present invention, and the amount of acicular ferrite is low, which deviates from the range of the present invention. Therefore, the amount of BH, ΔTS, and aging resistance at room temperature decrease.

B of steel plate No. 13 deviates from the appropriate range of the present invention, and the amount of acicular ferrite is low, which deviates from the range of the present invention, so the r value, BH amount, ΔTS, and aging resistance at room temperature decrease. In addition, the Nb of steel sheet No. 14 deviates from the range of the present invention, and the amount of solid solution N is low, which deviates from the range of the present invention, so the strain age hardening property deteriorates.

N in steel plate No. 15 deviates from the suitable range of the present invention, there is little solid solution N, and the strain age hardening property decreases. The hot rolling conditions and cold rolling annealing conditions of steel sheets No. 17 to No. 20 deviate from the appropriate range, the microstructure is outside the scope of the present invention, the BH amount and ΔTS decrease, the strain age hardening characteristics decrease, and the aging resistance at room temperature deteriorates.

Example 5

Adopt the same method as embodiment 4 to make the steel of composition shown in table 12 into slab, heat this slab under the condition shown in table 13, carry out rough rolling to form the thin slab with 25mm thickness, then by carrying out the slab shown in table 13 Conditional finish rolling of hot rolling process to make hot rolled sheet. Furthermore, after the rough rolling, the front and rear adjacent thin slabs are joined to each other by fusion bonding on the side of the finish rolling, and continuous rolling is performed. In addition, the thin slab edge heater and the thin slab heater of the induction heating method are used to heat the widthwise end and the longitudinal end of the thin slab to adjust the temperature of the thin slab.

These hot-rolled sheets were made into cold-rolled sheets having a thickness of 1.6 mm through a cold-rolling process consisting of pickling and cold-rolling under the conditions shown in Table 13. Next, these cold-rolled sheets were subjected to continuous annealing under the conditions shown in Table 13.

Regarding the obtained cold-rolled annealed sheet, the solid solution N content, microstructure, tensile properties, r value measurement, and strain age hardening properties were investigated in the same manner as in Example 4. In addition, with regard to the width direction and the longitudinal direction of each cold-rolled annealed sheet, tensile properties were investigated at each of 10 locations, and variations in yield strength, tensile strength, and elongation were investigated.

These results are shown in Table 14.

The examples of the present invention all have excellent strain age hardening characteristics and high r values, and exhibit particularly high BH amounts, ΔTS, and average r values stably despite fluctuations in manufacturing conditions. In addition, it was confirmed that in the example of the present invention, through the implementation of continuous rolling and temperature adjustment in the longitudinal and width directions of the thin slab, the thickness accuracy and shape of the finished steel plate are improved, and the material dispersion is reduced to 1/2. In addition, the elongation of temper rolling was changed to 0.5 to 2%, and the elongation of leveling was changed to 0 to 1%, but the strain age hardening characteristics did not decrease.

Example 6

Embodiments of the present invention will be described below.

The molten steel with the composition shown in Table 15 was smelted in a converter and made into a steel slab by continuous casting. These steel slabs were heated (a part of the hot slabs were charged) under the conditions shown in Table 16, rough rolled into thin slabs, and then hot-rolled sheets were produced through a hot rolling process of performing finish rolling under the conditions shown in Table 16. In addition, some of the thin slabs are joined to each other by the fusion bonding method, and the front and rear adjacent thin slabs are joined together, and continuous rolling is performed.

These hot-rolled sheets were made into cold-rolled sheets through a cold-rolling process consisting of pickling and cold-rolling under the conditions shown in Table 16. Next, under the conditions shown in Table 16, these cold-rolled steel sheets were subjected to box annealing and continuous annealing subsequent to the box annealing. Some of them are subjected to temper rolling after the cold rolling and annealing process. Furthermore, the case of unboxing annealing was also carried out. The annealing temperature of the box annealing is all made above the recrystallization temperature.

Regarding the obtained cold-rolled annealed sheet, the solid solution N content, microstructure, tensile properties, r-value measurement, strain age hardening properties, and aging properties were investigated.

In addition, after the continuous annealing in the table, hot-dip galvanizing was performed on the surface of the steel sheets of No. 17 and No. 18 in the production line to produce coated steel sheets, and various characteristics were evaluated in the same manner.

These results are shown in Table 17.

In the examples of the present invention, all exhibited excellent ductility, exceptionally high BH content and ΔTS, excellent strain age hardening properties, a high r value with an average r value of 1.2 or more, and non-aging properties at room temperature. In addition, the characteristics of the hot-dip galvanized steel sheets of steel sheets No. 17 and No. 18 shown in Table 17 are similarly similar to those of the continuously annealed cold-rolled steel sheets. On the contrary, the comparative examples deviated from the scope of the present invention had poor ductility, little BH amount, ΔTS, or significant aging deterioration, and did not have all the target properties, and could not be said to be steel sheets with sufficient properties.

In steel plate No. 11, the C and N amounts deviate from the range of the present invention, and the solid solution N amount and martensite amount are lower than the range of the present invention, so the BH amount and ΔTS decrease, and ΔY-E1 increases. In addition, Al, N/Al, and N of steel plate No. 12 deviate from the scope of the present invention, the amount of solid solution N is low, and deviate from the scope of the present invention, and the average grain size of ferrite is large, and deviates from the scope of the present invention. Therefore, The amount of BH and ΔTS decreased, and ΔY-E1 increased.

The slab heating temperature and FDT of steel plate No. 13 deviate from the appropriate range of the present invention, the amount of solid solution N and the amount of martensite are lower, and deviate from the scope of the present invention, and the average grain size of ferrite is large, which deviates from the scope of the present invention. Therefore, the r value, BH amount, and ΔTS decrease. In addition, the coiling temperature of steel plate No. 14 after hot rolling deviates from the scope of the present invention, and the amount of solid solution N is low, which deviates from the scope of the present invention. The ferrite The average crystal grain size is large and deviates from the range of the present invention, so the r value, BH amount, and ΔTS are reduced.

The continuous annealing temperature of steel plate No. 15 deviates from the suitable range of the present invention, martensite cannot be formed, and the average grain size of ferrite is large, which deviates from the range of the present invention. Therefore, the BH amount and ΔTS decrease, and ΔY-E1 increases. In addition, steel plate No. 16 was not subjected to box annealing, and the desired texture was not developed, so the r value was particularly low. In addition, the average grain size of ferrite and the area ratio of martensite also deviate from the range of the present invention.

Example 7

Using the same method as in Example 1, the steel of the composition shown in Table 18 was made into a slab, the slab was heated under the conditions shown in Table 19, and rough rolling was carried out to form a thin slab with a thickness of 30 mm. The hot rolling process of the finish rolling under the indicated conditions is made into a hot rolled sheet. In addition, some of them were continuously rolled by joining front and rear adjacent thin slabs to each other on the side of the finish rolling after the rough rolling by fusion bonding. In addition, the thin slab edge heater and the thin slab heater of the induction heating method are used to heat the widthwise end and the longitudinal end of the thin slab to adjust the temperature of the thin slab.

These hot-rolled sheets were made into 1.6 mm-thick cold-rolled sheets through a cold-rolling process consisting of pickling and cold-rolling under the conditions shown in Table 19. Next, these cold-rolled sheets were subjected to box annealing under the conditions shown in Table 19, and then to continuous annealing using a continuous annealing furnace. Furthermore, the annealing temperature of the box annealing is set to be above the recrystallization temperature.

The obtained cold-rolled annealed sheet was investigated in the same manner as in Example 1, for the solid solution N content, microstructure, tensile properties, r value measurement, and strain age hardening properties. In addition, with regard to the width direction and the longitudinal direction of each cold-rolled annealed sheet, tensile properties were investigated at each of 10 locations, and variations in yield strength, tensile strength, and elongation were investigated. Note that dispersion is represented by the difference between the maximum value and the minimum value at all the sites measured, for example, δYS=(maximum value of YS)−(minimum value of YS). The results are shown in Table 20.

The examples of the present invention all have excellent strain age hardening properties and high r-values, and exhibit exceptionally high BH amounts, ΔTS, and average r-values stably despite fluctuations in manufacturing conditions. In addition, it was confirmed that in the example of the present invention, by implementing continuous rolling and temperature adjustment in the longitudinal and width directions of the thin slab, the thickness accuracy and shape of the product steel plate are improved, and the material dispersion is reduced.

Industrial Utilization Possibility

According to the present invention, it is possible to obtain a cold-rolled steel sheet in which TS is greatly increased by press forming-heat treatment while ensuring excellent deep drawability during press forming. And the following excellent effects can be obtained: the cold-rolled steel sheet can be used to industrially manufacture electro-galvanized steel sheets, hot-dip galvanized steel sheets, and alloyed hot-dip galvanized steel sheets. Table 1 steel Composition Composition (mass%) (1)'Formula ^* (2)' Formula ^** Remark C N Si mn B al Nb P S A 0.0009 0.011 0.01 0.12 0.0009 0.010 0.016 0.009 0.005 0.0023 -0.0001 Suitable example B 0.0020 0.015 0.01 0.11 0.0011 0.012 0.035 0.015 0.006 0.0021 -0.0003 " C 0.0005 0.009 0.01 0.09 0.0005 0.009 0.010 0.011 0.005 0.0022 -0.0001 " D. 0.0020 0.021 0.01 0.50 0.0015 0.020 0.035 0.030 0.004 0.0035 -0.0003 " E. 0.0003 0.010 0.50 0.12 0.0006 0.011 0.099 0.045 0.010 0.0022 -0.0003 " f 0.0011 0.030 0.80 0.80 0.0011 0.028 0.025 0.009 0.005 0.0103 -0.0005 " G 0.0011 0.018 0.70 0.12 0.0008 0.012 0.018 0.008 0.005 0.0080 -0.0001 " h 0.0005 0.020 0.35 0.11 0.0014 0.020 0.025 0.007 0.005 0.0041 -0.0011 " I 0.0098 0.002 0.01 0.12 0.0007 0.038 0.055 0.009 0.005 -0.0269 0.0027 comparative example J 0.0022 0.001 0.50 0.12 0.0008 0.012 0.001 0.008 0.005 -0.0064 0.0021 " K 0.0260 0.003 0.02 0.25 0.0001 0.035 0.001 0.013 0.007 -0.0154 0.0259 " L 0.0027 0.011 0.01 0.12 0.0007 0.014 0.001 0.009 0.005 0.0027 0.0026 "

* (1)' formula: N%-(14/93 Nb%+14/27 Al%+14/11 B%) (0.0015 or more is the suitable range of the present invention)

**(2)' formula: C%-(0.5 12/93 Nb%) (below 0 is the suitable scope of the present invention) table 2 No. steel Hot rolling condition Cold rolling process conditions Remark Slab heating temperature (℃) Finish rolling finish temperature (°C) Cooling conditions after finish rolling (s, ℃/s) Winding temperature (℃) Cold rolling reduction (%) Heating rate (℃/s) Recrystallization annealing temperature (°C) With or without alloying hot-dip galvanized treatment 1 A 1150 920 0.39, 58 600 80 15 840 have Suitable example 2 B 1100 900 0.32, 53 550 80 15 840 " " 3 C 1130 650 ^* 0.25, 15 600 80 15 840 none " 4 D. 1110 900 0.35, 59 400 80 15 840 " " 5 E. 1160 920 0.28,70 650 80 15 840 have " 6 f 1150 910 0.39, 65 550 80 15 840 " " 7 G 1150 900 0.36, 58 500 80 15 840 none " 8 h 1000 680 ^* 0.26, 20 500 80 15 840 " " 9 A 1150 920 1.13, 10 600 80 25 840 have " 10 E. 1160 920 0.67, 15 650 80 25 840 " " 11 I 1140 920 0.35, 57 650 80 15 840 have comparative example 12 J 1000 900 0.68,7 550 80 15 840 " " 13 K 1150 880 0.38, 61 500 80 15 840 none " 14 L 1000 920 1.02, 9 600 80 15 940 " " ^*Finish rolling finish temperature is less than Ar₃phase transition point. The cooling conditions after finish rolling represent the cooling start time (s) and the cooling rate (°C/s). table 3 No. steel Product characteristics Forming - change in tensile strength after heat treatment Remark Tensile strengthTS(MPa) r value Grain size (μm) ΔE1(%) TS(MPa) after heat treatment at 120℃ ΔT.S.(MPa) after heat treatment at 120℃ TS(MPa) after heat treatment at 170℃ ΔT.S.(MPa) after heat treatment at 170℃ 1 A 360 1.8 18 1.6 435 75 455 95 Suitable example 2 B 365 1.7 12 1.2 460 95 470 105 " 3 C 355 2.3 19 1.8 420 65 440 85 " 4 D. 390 1.9 14 1.3 480 90 490 100 " 5 E. 430 1.8 18 1.7 520 90 530 100 " 6 f 460 1.6 18 1.7 575 115 595 135 " 7 G 430 1.8 13 1.4 540 110 555 125 " 8 h 390 2.1 19 1.7 490 100 510 120 " 9 A 352 1.9 twenty two 2.4 414 62 440 88 " 10 E. 417 1.9 twenty three 2.7 489 72 503 86 " 11 I 370 1.9 18 3.2 420 50 450 80 comparative example 12 J 400 1.6 twenty two 2.4 435 35 460 60 " 13 K 360 1.1 26 3.1 390 30 415 55 " 14 L 360 1.2 25 4.2 414 54 445 85 " Table 4table 5 Al% N/Al TS×r value MPa ΔTSMPa 0.020 0.75 775 58 0.036 0.42 762 55 0.049 0.31 753 42 0.072 0.21 720 25 0.080 0.19 719 19 Table 6 steel C% Si% Mn% P% S% N% Al% Nb% Ti% V% B% N/(Al+Nb+Ti+V+B)* Ar ₃ ℃ A 0.0013 0.01 0.15 0.009 0.005 0.0135 0.011 --- --- --- --- 1.23 884 B 0.0012 0.50 0.50 0.011 0.004 0.0175 0.010 0.010 --- --- 0.0010 0.83 888 C 0.0008 0.01 0.01 0.009 0.003 0.0190 0.009 --- 0.015 0.015 0.0010 0.48 896 D. 0.0025 0.01 0.09 0.005 0.003 0.0160 0.010 0.010 0.015 --- 0.0006 0.45 876 E. 0.0016 0.55 0.80 0.04 0.003 0.0195 0.015 0.009 --- 0.018 0.0012 0.45 905 f 0.0011 0.75 0.75 0.02 0.002 0.0135 0.008 0.015 --- --- 0.0010 0.56 889 G 0.0032 0.01 0.14 0.009 0.003 0.0130 0.045 --- --- --- --- 0.29 894 h 0.0019 0.45 0.45 0.04 0.005 0.0140 0.015 0.035 --- --- 0.0012 0.27 882 *For non-added Nb, Ti, V, B, record the above concentration as "---", and take the concentration as "0" when calculating N/(Al+Nb+Ti+V+B). Table 7 No. steel hot rolled Hot rolled sheet annealing cold rolling Cold rolled sheet recrystallization annealing Cooling: until the cooling rate before coating ℃/s coating alloying Remark rough rolling Finish rolling Reduction rate% wind up temperature °C Process intermittent continuous Plate temperature ℃ Bath temperature °C temperature °C Hold time s CT°C SRT°C RDT°C FET°C With or without lubrication FDT°C Reduction rate% temperature °C time s 1 A 1150 900 810 have 660 95 510 790 Intermittent 80.0 860 40 15 470 465 - - II 2 B 1150 910 830 have 660 95 510 840 continuous 80.0 860 40 - - - - - I 3 C 1140 920 840 have 680 95 520 800 Intermittent 82.5 870 40 30 465 460 460 25 III 4 D. 1180 900 820 have 650 95 500 830 continuous 82.5 850 40 35 470 465 - - II 5 E. 1190 930 850 have 690 95 530 810 Intermittent 80.0 880 40 25 465 460 460 twenty three III 6 f 1160 920 840 have 670 95 530 810 Intermittent 82.5 880 40 - - - - - I 7 G 1120 910 820 have 670 95 520 800 Intermittent 80.0 870 40 50 480 470 - - II 8 h 1100 900 810 have 650 95 500 790 Intermittent 80.0 860 40 15 465 470 470 20 III 9 A 1080 900 810 have 660 95 500 790 Intermittent 82.5 860 40 15 470 465 - - II 10 B 1140 910 820 have 670 65 520 790 Intermittent 82.5 860 40 1 470 460 460 20 III 11 C 1250 1150 1070 have 910 95 750 850 continuous 80.0 870 40 30 465 460 460 25 III 12 D. 1170 900 810 have 660 95 510 not dealt with 80.0 850 40 - - - - - I SRT = slab heating temperature, RDT = rough rolling exit temperature, FET = finish rolling entry temperature, FDT = finish rolling exit temperature, CT = coiling temperature I = cold rolled steel plate, II = hot dip galvanized steel plate, III = alloyed hot-dip galvanized steel sheet Table 8 No. steel Solid solution N% Tensile properties before strain aging Strain age hardening properties Remark YSMPa TSMPa E1% r value TSxr value MPa BHMPa ΔTS MPa 1 A 0.0069 225 321 53 2.4 770 122 75 Example 2 B 0.0089 274 391 43 2.3 899 183 93 Example 3 C 0.0054 221 316 54 2.8 885 97 72 Example 4 D. 0.0049 221 316 54 2.8 885 90 66 Example 5 E. 0.0050 304 435 39 2.0 870 80 63 Example 6 f 0.0088 304 434 39 2.1 911 133 87 Example 7 G 0.0000 224 320 53 2.8 896 3 0 comparative example 8 h 0.0000 284 405 42 2.1 851 2 0 comparative example 9 A 0.0070 215 311 50 2.3 715 152 83 comparative example 10 B 0.0082 274 391 43 1.9 743 143 94 comparative example 11 C 0.0035 236 331 51 2.0 662 93 48 comparative example 12 D. 0.0041 241 336 51 2.1 706 85 49 comparative example Table 9 steel Chemical composition (mass%) C Si mn P S Al N N/Al Nb 12/93 Nb B other A 0.0025 0.15 0.85 0.050 0.002 0.005 0.0120 2.40 0.0040 0.0005 0.0015 - B 0.0050 0.51 1.20 0.004 0.001 0.008 0.0150 1.88 0.0070 0.0009 0.0009 - C 0.0024 0.15 0.35 0.040 0.001 0.004 0.0120 3.00 0.0050 0.0006 0.0015 - D. 0.0150 0.15 0.88 0.010 0.002 0.010 0.0100 1.00 0.0102 0.0013 0.0012 - E. 0.0023 0.01 1.40 0.005 0.002 0.011 0.0120 1.09 0.0070 0.0009 0.0010 - f 0.0023 0.05 1.35 0.045 0.001 0.007 0.0120 1.71 0.0080 0.0010 0.0020 Mo: 0.15 G 0.0025 0.15 1.25 0.007 0.001 0.004 0.0110 2.75 0.0070 0.0009 0.0015 Ti: 0.013 h 0.0025 0.25 1.25 0.008 0.001 0.011 0.0140 1.27 0.0060 0.0008 0.0014 Cu: 0.50, Ni: 0.20 I 0.0024 0.15 1.25 0.005 0.003 0.011 0.0150 1.36 0.0089 0.0010 0.0018 Ni: 0.05, V: 0.02 J 0.022 0.15 1.21 0.008 0.002 0.005 0.0120 2.40 0.0076 0.0011 0.0020 Cu: 0.10, Ni: 0.05 K 0.070 0.25 1.50 0.015 0.003 0.055 0.0040 0.07 0.0090 0.0012 0.0021 - L 0.0025 0.15 0.85 0.050 0.002 0.005 0.0120 2.40 0.0001 0.0000 ≤0.0001 - m 0.0025 0.15 0.85 0.050 0.002 0.005 0.0120 2.40 0.0050 0.0006 ≤0.0001 - o 0.0025 0.17 0.81 0.008 0.001 0.007 0.0120 1.71 0.110 0.0142 0.0015 - P 0.0026 0.17 0.82 0.007 0.001 0.007 0.0400 5.71 0.0070 0.0009 0.0017 - Q 0.0025 0.18 0.81 0.008 0.001 0.007 0.0120 1.71 0.0080 0.0010 0.0022 Ca: 0.0035 Table 10 Steel plate No. steel hot rolling process Cold rolling process α-γ coexistence temperature zone ℃ Cold rolled sheet annealing process Quenched and tempered rolling Remark continuous annealing Slab Heating Temperature SRT℃ rough rolling Whether the thin slabs are joined or not Finish rolling wind up Cold rolling reduction % Cold-rolled plate thickness mm Thin slab thickness mm Ar ₃ phase transition point ℃ Exit temperature FDT°C Thickness of hot rolled plate mm Winding temperature CT°C Annealing temperature ℃ Hold time s Cooling rate ℃/s Cooling stop temperature ℃ type Elongation% 1 A 1210 28 none 873 900 4.0 550 83 0.70 860～950 880 35 30 350 Skin-pass rolling 0.5 Example of the invention 2 B 1230 28 none 837 900 4.0 530 81 0.75 835～935 870 30 35 350 Skin-pass rolling 0.5 Example of the invention 3 C 1220 28 none 908 910 4.0 550 83 0.70 865～955 885 25 25 350 Skin-pass rolling 0.5 Example of the invention 4 D. 1180 28 none 851 890 4.0 530 83 0.70 845～922 885 40 45 350 - - Example of the invention 5 E. 1190 28 none 801 890 4.0 520 83 0.70 825～906 875 25 40 450 leveling 0.5 Example of the invention 6 f 1180 25 none 827 890 4.5 500 78 1.00 855～930 885 30 35 450 Skin pass cold rolling + leveling 0.5 Example of the invention 7 G 1180 25 none 820 890 6.0 570 80 1.20 835～918 875 45 30 450 Skin pass cold rolling + leveling 0.7 Example of the invention 8 h 1180 25 none 824 890 4.0 610 80 0.80 835～923 880 25 45 450 Skin pass cold rolling + leveling 0.7 Example of the invention 9 I 1190 28 none 819 890 4.0 550 84 0.65 830～915 865 25 30 300 Skin pass cold rolling + leveling 0.7 Example of the invention 10 J 1210 28 none 823 870 4.0 450 83 0.70 830～910 870 25 30 300 Skin pass cold rolling + leveling 0.7 Example of the invention 11 K 1200 28 none 808 880 4.0 520 83 0.70 755～895 870 25 30 300 Skin pass cold rolling + leveling 0.5 comparative example 12 L 1200 28 none 873 900 4.0 550 83 0.70 860～950 880 35 30 300 Skin pass cold rolling + leveling 0.5 comparative example 13 m 1200 25 none 873 900 4.0 550 83 0.70 860～950 880 35 30 300 Skin pass cold rolling + leveling 0.5 comparative example 14 o 1210 25 none 856 890 4.0 520 83 0.70 850～940 890 30 35 300 Skin pass cold rolling + leveling 0.5 comparative example 15 P 1240 25 none 854 880 4.0 520 83 0.70 850～940 890 30 35 300 Skin-pass rolling 0.7 comparative example 16 Q 1220 25 none 856 900 4.0 530 83 0.70 850～940 880 30 35 300 Skin-pass rolling 0.7 Example of the invention 17 A 950 25 none 870 750 4.0 540 83 0.70 855～950 880 30 30 300 Skin-pass rolling 0.7 comparative example 18 A 1200 28 none 873 900 4.0 780 83 0.70 865～950 870 35 35 300 Skin-pass rolling 0.7 comparative example 19 A 1190 28 none 873 905 4.0 540 83 0.70 860～950 820 25 40 300 Skin-pass rolling 0.7 comparative example 20 A 1190 28 none 873 890 4.0 550 83 0.70 860～950 870 30 5 300 Skin-pass rolling 0.7 comparative example Table 11 Steel plate No. steel Steel plate solid solution N amount mass% steel structure Product board characteristics Strain age hardening properties Aging resistance Remark ferrite Phase 2 tensile properties BH content MPa ΔTS MPa ΔY-E1% Area ratio% Particle size μm AF area ratio% YSMPa TSMPa E1% mean r value 1 A 0.0090 92 7 8 315 448 38 1.8 85 75 0.0 Example of the invention 2 B 0.0100 95 8 5 355 510 34 1.7 90 80 0.0 Example of the invention 3 C 0.0085 92 7 8 295 420 41 1.8 80 75 0.0 Example of the invention 4 D. 0.0095 92 6 8 334 475 36 1.4 95 80 0.0 Example of the invention ^** 5 E. 0.0100 93 7 7 325 465 37 1.7 88 75 0.0 Example of the invention 6 f 0.0105 90 7 10 385 550 35 1.5 102 75 0.0 Example of the invention 7 G 0.0105 88 6 12 315 455 39 1.8 88 72 0.0 Example of the invention 8 h 0.0095 92 10 8 325 460 37 1.8 98 75 0.0 Example of the invention 9 I 0.0090 90 5 10 320 465 37 1.8 97 85 0.0 Example of the invention 10 J 0.0105 88 5 12 420 595 29 1.4 105 85 0.0 Example of the invention ^** 11 K 0.0002 75 7 25 370 545 25 1.1 45 34 1.5 comparative example 12 L 0.0080 98 15 2 290 420 35 1.3 15 10 1.5 comparative example 13 m 0.0085 97 8 3 290 410 33 1.2 15 7 1.0 comparative example 14 o 0.0003 98 6 2 320 455 33 1.1 5 15 0.8 comparative example 15 P 0.0190 95 7 5 310 440 37 1.6 90 70 2.8 comparative example 16 Q 0.0098 92 7 8 315 450 39 1.7 90 70 0 Example of the invention 17 A 0.0009 98 twenty two 2 275 420 33 1.1 0 15 0.5 comparative example 18 0.0040 95 25 5 275 410 34 1.3 30 10 0.5 comparative example 19 0.0050 100 12 0 270 380 33 1.3 25 12 0.8 comparative example 20 0.0030 99 twenty two 1 265 385 31 1.3 20 10 0.9 comparative example

AF: acicular ferrite

^** : With hot-dip galvanized treatment Sheet 12 steel Chemical composition (mass%) C Si mn P S Al N N/Al Nb 12/93Nb B other N 0.0085 0.005 0.55 0.009 0.005 0.015 0.0126 0.84 0.008 0.0010 0.0015 - Table 13 Steel plate No. steel hot rolling process Cold rolling process α-γ coexistence temperature zone ℃ Cold rolled sheet annealing process modulation rolling Remark continuous annealing Slab Heating Temperature SRT℃ rough rolling Whether the thin slabs are joined or not Finish rolling wind up Cold reduction % Cold-rolled plate thickness mm Thin slab thickness mm Ar ₃ phase transition point ℃ Material temperature FDT°C Thickness of hot rolled plate mm Winding temperature CT°C Annealing temperature ℃ Hold time s Cooling rate ℃/s Cooling stop temperature ℃ type Elongation% 2-1 N 1250 25 have ^* 870 890 4.5 570 69 1.4 840～900 880 30 45 300 Skin-pass rolling 0.5 Example of the invention 2-2 1230 25 have ^* 870 890 4.5 550 69 1.4 840～900 880 35 45 310 Skin-pass rolling 0.5 Example of the invention 2-3 1240 25 have ^** 870 890 4.5 560 69 1.4 840～900 880 30 45 310 Skin-pass rolling 0.5 Example of the invention

^* ) Implement lubricated rolling

^** ) Lubricated rolling is carried out, using thin slab heaters and slab edge heaters Table 14 Steel plate No. steel Steel plate solid solution N amount mass% steel structure Product board characteristics Product board characteristics Strain age hardening properties Aging resistance Remark ferrite Phase 2 tensile properties Discretization of Tensile Properties* BH content MPa ΔTS MPa ΔY-E1% Area ratio% Particle size μm AF area ratio% YSMPa TSMPa E1% mean r value δYSMPa δTSMPa δE1% 2-1 N 0.0120 91 7 9 315 445 37 1.6 10 10 5 15 15 0 Example of the invention 2-2 N 0.0120 91 7 9 319 447 37 1.6 5 5 2 7 6 0 Example of the invention 2-3 N 0.0125 92 7 8 318 450 38 1.5 3 5 1 8 5 0 Example of the invention

^* ) δYS, δTS, δE1: = (maximum value - minimum value)

^** ) AF: acicular ferrite, M: martensite, B: bainite, P: pearlite Table 15 steel Chemical composition (mass%) Phase transition point (°C) C Si mn P S al N N/Al other Ac ₁ Ac ₃ A 0.050 0.01 0.55 0.040 0.001 0.004 0.0120 3.00 - 735 858 B 0.080 0.02 0.30 0.060 0.001 0.008 0.0153 1.91 - 749 866 C 0.070 0.15 0.45 0.040 0.001 0.002 0.0118 5.90 REM: 0.0015 740 862 D. 0.055 0.15 0.88 0.050 0.002 0.006 0.0095 1.58 Ca: 0.0020 739 863 E. 0.030 0.01 1.40 0.005 0.002 0.011 0.0120 1.09 Ca: 0.0025, REM: 0.0020 706 832 f 0.050 0.05 0.55 0.045 0.001 0.007 0.0123 1.76 Mo: 0.02 739 862 G 0.065 0.02 0.75 0.045 0.001 0.004 0.0114 2.85 Ti: 0.013, B: 0.0005 735 855 h 0.050 0.01 0.85 0.008 0.001 0.008 0.0138 1.73 Cu: 0.50, Ni: 0.20, Cr: 0.20 717 834 I 0.040 0.01 0.55 0.005 0.003 0.011 0.0151 1.37 Ni: 0.05, V: 0.02 717 842 J 0.120 0.15 0.95 0.008 0.002 0.005 0.0118 2.36 Cu: 0.10, Ni: 0.05, Nb: 0.007 719 830 K 0.010 0.25 0.96 0.015 0.003 0.002 0.0042 2.10 - 723 857 L 0.080 0.15 0.85 0.050 0.002 0.045 0.0040 0.09 - 739 859 Table 16 Steel plate No. steel hot rolling process Cold rolling process Cold rolled sheet annealing process Quenched and tempered rolling Slab Heating Temperature SRT℃ rough rolling Whether the thin slabs are joined or not Finish rolling Cooling after cold rolling wind up Cold reduction % Cold-rolled plate thickness mm Box annealing continuous annealing Overdue Thin slab thickness mm Material temperature FDT°C Thickness of hot rolled plate mm Start time Δts Cooling rate V°C/s Winding temperature CT°C Annealing temperature ℃ Annealing temperature ℃ Hold time s Cooling rate ℃/s Cooling stop temperature ℃ Residence time in the temperature zone above 350°C ^** s Elongation% 1 A 1210 30 none 850 4.0 0.2 50 550 83 0.70 700 790 40 45 270 - 0.2 2 B 1230 30 none 870 4.0 0.2 50 530 81 0.75 680 780 30 50 270 - 0.2 3 C 1220 30 none 840 4.0 0.1 50 550 83 0.70 700 800 30 55 270 - 0.2 4 D. 1180 30 none 850 4.0 0.5 50 530 83 0.70 690 800 45 45 250 - 0.2 5 E. 1190 30 none 850 4.0 0.2 45 520 83 0.70 720 770 25 45 200 - 0.5 6 f 1180 ^*** 35 none 864 4.5 0.2 45 500 78 1.00 700 780 35 55 200 - - 7 G 1180 35 none 860 6.0 1.2 45 570 80 1.20 700 830 50 60 300 - 0.5 8 h 1180 35 none 830 4.0 0.3 45 610 80 0.80 700 810 25 50 400 50 0.5 9 I 1190 30 have 840 4.0 ^* 0.3 45 550 84 0.65 700 790 25 45 250 - 0.5 10 J 1210 30 none 840 4.0 0.3 50 450 83 0.70 700 770 25 30 250 - 0.5 11 K 1200 30 none 850 4.0 0.3 50 520 83 0.70 700 800 25 30 250 - 0.5 12 L 1200 30 none 880 4.0 0.2 50 550 83 0.70 700 790 35 30 270 - 0.7 13 A 950 35 none 750 4.0 0.5 50 540 83 0.70 700 800 30 30 270 - 0.7 14 A 1200 35 none 880 4.0 0.3 50 780 83 0.70 700 830 35 35 250 - 0.7 15 A 1190 35 none 870 4.0 0.2 50 540 83 0.70 700 680 25 40 250 - 0.5 16 A 1200 35 none 880 4.0 0.2 50 550 83 0.70 - 820 29 35 250 - 0.5 17 B 1190 30 none 840 4.0 0.3 50 520 83 0.70 710 780 35 70 450 - ^**** 18 G 1180 30 none 850 4.0 0.3 50 520 83 0.70 720 780 40 60 450 - ^****

^* ) Implement lubricated rolling

^** ) Residence time below the cooling stop temperature and above 350°C ^*** ) Loading of hot slabs ^**** ) Tempering rolling with an elongation of 0.5% after hot-dip galvanizing Table 17 Steel plate No. steel Steel plate solid solution N amount mass% steel structure Product board characteristics Strain age hardening properties Aging resistance Remark ferrite Phase 2 tensile properties BH content MPa ΔTSMPa ΔY-E1% Area ratio% Particle size μm M area ratio% Other Phases: Types YSMPa TSMPa E1% mean r value 1 A 0.0075 86 7 12 B 290 550 35 1.6 80 80 0.0 Example of the invention 2 B 0.0055 92 8 7 P 295 556 35 1.5 90 85 0.0 Example of the invention 3 C 0.0055 90 7 10 - 285 555 34 1.6 85 75 0.0 Example of the invention 4 D. 0.0075 89 6 10 B 295 565 34 1.6 95 80 0.0 Example of the invention 5 E. 0.0110 91 7 9 - 315 605 33 1.7 85 75 0.0 Example of the invention 6 f 0.0110 90 7 10 - 309 595 35 1.6 100 75 0.0 Example of the invention 7 G 0.0095 85 6 14 B 314 605 35 1.6 85 70 0.0 Example of the invention 8 h 0.0078 87 10 12 B 318 615 33 1.6 96 75 0.0 Example of the invention 9 I 0.0086 86 5 12 P 276 525 37 1.5 90 80 0.0 Example of the invention 10 J 0.0095 84 5 15 P 324 620 33 1.5 102 90 0.0 Example of the invention 11 K 0.0003 98 7 1 P 235 355 38 1.4 45 34 1.5 comparative example 12 L 0.0005 93 15 7 - 241 370 39 1.3 15 10 1.5 comparative example 13 A 0.0009 98 twenty two 0.5 P 285 430 32 1.1 0 15 0.5 comparative example 14 0.0009 95 25 5 - 285 445 32 1.2 30 10 0.5 comparative example 15 0.0010 100 12 0 - 315 480 34 1.4 25 12 0.8 comparative example 16 0.0020 99 twenty two 1 - 350 545 25 1.0 20 10 0.9 comparative example 17 B 0.0060 92 8 7 P 297 555 35 1.5 92 85 0 Example of the invention 18 G 0.0055 92 8 7 P 296 557 34 1.5 90 85 0 Example of the invention

M: Martensite, B: Bainite, P: Pearlite Table 18 steel Chemical composition (mass%) Phase transition point (°C) C Si mn P S Al N N/Al other Ac ₁ Ac ₃ m 0.052 0.01 0.60 0.035 0.001 0.002 0.0125 6.25 - 735 855 Table 19 Steel plate No. steel hot rolling process Cold rolling process Cold rolled sheet annealing process Quenched and tempered rolling Remark Slab Heating Temperature SRT℃ rough rolling Whether the thin slabs are joined or not Finish rolling Cooling after cold rolling wind up Cold reduction % Cold-rolled plate thickness mm Box annealing continuous annealing Overdue Thin slab thickness mm Material temperature FDT°C Thickness of hot rolled plate mm Start time Δts Cooling rate V°C/s Winding temperature CT°C Annealing temperature ℃ Annealing temperature ℃ Hold time s Cooling rate ℃/s Cooling stop temperature ℃ Residence time in the temperature range above 350°C s Elongation% 2-1 m 1180 30 none 850 4.5 0.3 45 540 64 1.6 700 790 40 45 270 - 0.2 Example of the invention 2-2 1200 30 have ^* 850 4.5 0.3 45 540 64 1.6 700 790 40 45 270 - 0.2 Example of the invention 2-3 1190 30 have ^** 855 4.5 0.3 45 540 64 1.6 700 700 40 45 270 - 0.2 Example of the invention

^* ) Implement lubricated rolling

^** ) Using thin slab heaters, slab edge heaters Table 20 Steel plate No. steel Steel plate solid solution N amount mass% steel structure Product board characteristics Product board characteristics Strain age hardening properties Aging resistance Remark ferrite Phase 2 tensile properties Discretization of Tensile Properties ^* BH content MPa ΔTS MPa ΔY-E1% Area ratio% Particle size μm M area ratio% Other Phases: Type ^** YSMPa TSMPa E1% mean r value δYSMPa δTSMPa δE1% 2-1 m 0.0075 85 7 12 B 295 551 35 1.6 25 15 2 80 80 0 Example of the invention 2-2 m 0.0075 85 7 12 B 295 552 35 1.6 20 10 1 80 80 0 Example of the invention 2-3 m 0.0075 85 7 12 B 292 553 35 1.6 15 8 1 80 80 0 Example of the invention

^* )δ YS, δTS, δE1: = (maximum value - minimum value)

^** )M: martensite, B: bainite, P: pearlite

Claims (25)

1. cold-rolled steel sheet that has excellent strain aging hardening properties, it is characterized in that: % represents with quality, have contain that C:0.15% is following, Si:1.0% following, Mn:2.0% is following, P:0.1% is following, S:0.02% is following, Al:0.005～0.030%, N:0.0050～0.0400% and N/Al:0.30 N above, the solid solution attitude are more than 0.0010%, remainder is by forming that Fe and unavoidable impurities constitute.

2. according to the cold-rolled steel sheet that has excellent strain aging hardening properties of claim 1 record, it is characterized in that: on the basis of aforementioned component, also contain the composition more than 1 group or 2 groups in the following a group～d group of representing with quality %, wherein,

A group: among Cu, Ni, Cr, the Mo more than a kind or 2 kinds, the total amount is below 1.0%;

B group: among Nb, Ti, the V more than a kind or 2 kinds, the total amount is below 0.1%;

C group: the B below 0.0030%;

D group: Ca, REM a kind or 2 kinds, the total amount is 0.0010～0.010%.

3. cold-rolled steel sheet that has excellent strain aging hardening properties, it is characterized in that: % represents with quality, has to contain C: less than 0.01%, Si:0.005～1.0%, Mn:0.01～1.5%, P:0.1% are following, S:0.01% is following, Al:0.005～0.030%, N:0.005～0.040% and N/Al:0.30 N above, the solid solution attitude are more than 0.0010%, remainder is by forming that Fe and unavoidable impurities constitute.

4. according to the cold-rolled steel sheet that has excellent strain aging hardening properties of claim 3 record, it is characterized in that: on the basis of above-mentioned composition, also contain B:0.0001～0.0030%, Nb:0.005～0.050% of representing with quality % in the scope that satisfies following formula (1), (2), remainder is essentially Fe.

N％≥0.0015+14/93·Nb％+14/27·Al％+14/11·B％??---(1)

C％≤0.5·(12/93)·Nb％????????????????????????????---(2)

5. according to claim 3 or 4 cold-rolled steel sheets of putting down in writing that have excellent strain aging hardening properties, it is characterized in that: on above-mentioned composition basis, also containing the total amount of representing with quality % as required is Cu, the Ni below 1.0%, the composition more than a kind or 2 kinds among the Mo.

6. according to the cold-rolled steel sheet that has excellent strain aging hardening properties of claim 1～5 record, it is characterized in that: the size of microcrystal of steel plate is below the 20 μ m.

7. according to the cold-rolled steel sheet that has excellent strain aging hardening properties of claim 1～6 record, it is characterized in that: in thermal treatment temp: 120～200 ℃ cold zone has the intensity ascending amount after the shaping: more than the 60MPa.

8. an electro-galvanizing that has excellent strain aging hardening properties, galvanizing and alloyed hot-dip galvanized steel plate is characterized in that: the surface at the cold-rolled steel sheet of claim 1～7 record possesses electro-galvanizing, galvanizing and alloyed hot-dip zinc-coated layer.

9. the manufacture method of a cold-rolled steel sheet that has excellent strain aging hardening properties, it is characterized in that: % represents with quality, to contain C: less than 0.01%, Si:0.005～1.0%, Mn:0.01～1.5%, below the P:0.1%, below the S:0.01%, Al:0.005～0.030%, N:0.005～0.040%, and content range satisfies more than the N/Al:0.30, the steel billet that remainder is essentially the composition of Fe carries out hot rolling, at this moment, after finishing, finish rolling begins to cool down immediately, in the coiling temperature: reel down for 400～800 ℃, implement draft then: 60～95% cold rolling after, under 650～900 ℃ temperature, implement recrystallization annealing.

10. according to the manufacture method of the cold-rolled steel sheet that has excellent strain aging hardening properties of claim 9 record, it is characterized in that: on the basis of above-mentioned composition, also contain B:0.0001～0.0030%, Nb:0.005～0.050% of representing with quality % in the scope that satisfies following formula (1), (2), remainder is essentially Fe.

N％≥0.0015+14/93·Nb％+14/27·Al％+14/11·B％??---(1)

C％≤0.5·(12/93)·Nb％????????????????????????????---(2)

11. manufacture method according to claim 9 or 10 cold-rolled steel sheets of putting down in writing that have excellent strain aging hardening properties, it is characterized in that: in the temperature-rise period of above-mentioned recrystallization annealing, the humidity province of the 500 ℃～recrystallization temperature speed with 1～20 ℃/second is heated up.

12. the manufacture method of an alloy hot-dip galvanized steel sheet that has excellent strain aging hardening properties is characterized in that: in claim 9～11, after the recrystallization annealing, implement galvanizing and handle, then implement the heating Alloying Treatment.

13. deep-drawing cold-rolled steel sheet that has excellent strain aging hardening properties, it is characterized in that: % represents with quality, have and contain C: less than 0.01%, Si:0.005～1.0%, Mn:0.01～1.5%, P:0.1% are following, S:0.01% is following, Al:0.005～0.030%, N:0.005～0.040% and N/Al:0.30 N above, the solid solution attitude are more than 0.0010%, remainder is by forming that Fe and unavoidable impurities constitute, its TS * r value: more than the 750MPa.

14. the deep-drawing cold-rolled steel sheet that has excellent strain aging hardening properties according to claim 13 record, it is characterized in that: have on the basis of above-mentioned composition, also contain B:0.0001～0.0030%, Nb:0.005～0.050% of representing with quality % in the scope that satisfies following formula (1), (2), remainder is by forming that Fe and unavoidable impurities constitute, its TS * r value: more than the 750MPa.N％≥0.0015+14/93·Nb％+14/27·Al％+14/11·B％??---(1)C％≤0.5·(12/93)·Nb％????????????????????????????---(2)

15. the deep-drawing cold-rolled steel sheet that has excellent strain aging hardening properties according to claim 13 record, it is characterized in that: have on the basis that the steel of claim 13 record is formed, also contain B:0.0001～0.0030% of representing with quality %, Nb:0.005～0.050%, Ti:0.005～0.070%, the composition more than a kind or 2 kinds among V:0.005～0.10%, and N/ (Al+Nb+Ti+V+B): more than 0.30, the N of solid solution attitude is more than 0.0010%, remainder is by forming that Fe and unavoidable impurities constitute, its TS * r value: more than the 750MPa.

A 16. deep-drawing that the has excellent strain aging hardening properties manufacture method of cold-rolled steel sheet, it is characterized in that: % represents with quality, to have the C of containing: less than 0.01%, Si:0.005～1.0%, Mn:0.01～1.5%, below the P:0.1%, below the S:0.01%, Al:0.005～0.030%, N:0.005～0.040%, contain B:0.0001～0.0030%, Nb:0.005～0.050%, Ti:0.005～0.070%, composition more than a kind or 2 kinds among V:0.005～0.10%, and N/ (Al+Nb+Ti+V+B): after the steel billet material of the composition more than 0.30 was heated to more than 950 ℃, making roughing finish temperature was below 1000 ℃, Ar ₃More than, carry out roughing, then at Ar ₃Below, the humidity province more than 600 ℃ is while lubricating finish rolling, coiling, make that to begin to total draft that finish rolling is finished from roughing be more than 80% this moment, the hot-rolled sheet of gained is carried out recrystallization annealing, then carry out cold rollingly, the cold-reduced sheet of gained is carried out recrystallization annealing with draft 60～95%.

17. plasticity, strain-aged hardening characteristics and the good cold-rolled steel sheet of anti-room temperature ageing, it is characterized in that: % represents with quality, has the C:0.0015 of containing～0.025%, below the Si:1.0%, below the Mn:2.0%, below the P:0.1%, below the S:0.02%, below the Al:0.02%, N:0.0050～0.0250%, and contain B:0.0001～0.0050%, the composition more than a kind or 2 kinds of Nb:0.002～0.050%, and making N/Al is more than 0.3, the N of solid solution attitude is more than 0.0010%, remainder is by forming of constituting of Fe and unavoidable impurities and by the acicular ferrite of representing with area occupation ratio more than 5% mutually and median size: the tissue that the following ferritic phase of 20 μ m constitutes, its r value: more than 1.2.

18. the cold-rolled steel sheet according to claim 17 record is characterized in that: on the basis of aforementioned component, also contain the composition 1 group or 2 group or more of the following a group～c that represents with quality % among organizing, wherein,

A group: among Cu, Ni, Cr, the Mo more than a kind or 2 kinds, the total amount is below 1.0%;

B group: among Ti, the V a kind or 2 kinds, the total amount is below 0.1%;

C group: Ca, REM a kind or 2 kinds, the total amount is 0.0010～0.010%.

19. one kind has the r value: more than 1.2, plasticity, the manufacture method of the cold-rolled steel sheet that strain-aged hardening characteristics and anti-room temperature ageing are good, it is characterized in that: % represents with quality, to contain C:0.0015～0.025%, below the Si:1.0%, below the Mn:2.0%, below the P:0.1%, below the S:0.02%, below the Al:0.02%, N:0.0050～0.0250%, and contain B:0.0001～0.0050%, more than a kind or 2 kinds of Nb:0.002～0.050%, and N/Al is that the plate slab of the composition more than 0.3 is heated to slab heating temperature: more than 1000 ℃, carry out the roughing base that laminates, implement successively this thin slab execution finish rolling is gone out the material temperature: the finish rolling more than 800 ℃, and in the coiling temperature: make the hot-rolled process of hot-rolled sheet thereby reel below 800 ℃, thereby this hot-rolled sheet is implemented pickling and the cold rolling cold rolling process of making cold-reduced sheet, under the temperature in the ferritic-austenitic two-phase region this cold-reduced sheet is carried out continuous annealing, with speed of cooling: 10～300 ℃ of/second cold-reduced sheet annealing operations that are cooled to the humidity province below 500 ℃.

20. the manufacture method according to the cold-rolled steel sheet of claim 19 record is characterized in that: on the basis of aforementioned component, also contain the composition 1 group or 2 group or more of the following a group～c that represents with quality % among organizing, wherein,

A group: among Cu, Ni, Cr, the Mo more than a kind or 2 kinds, the total amount is below 1.0%;

B group: among Ti, the V more than a kind or 2 kinds, the total amount is below 0.1%;

C group: Ca, REM a kind or 2 kinds, the total amount is 0.0010～0.010%.

21. one kind has high r value and good strain-aged hardening characteristics and the non-ageing high strength cold rolled steel plate of normal temperature, it is characterized in that: % represents with quality, has the C:0.025 of containing～0.15%, below the Si:1.0%, below the Mn:2.0%, below the P:0.08%, below the S:0.02%, below the Al:0.02%, N:0.0050～0.0250%, and N/Al is more than 0.3, the N that contains the solid solution attitude more than 0.0010%, remainder is by forming that Fe and unavoidable impurities constitute, with contain the average crystal grain particle diameter of representing with area occupation ratio more than 80%: the ferritic phase that 10 μ m are following, also contain tissue, its r value: more than 1.2 as the martensitic phase of representing with area occupation ratio more than 2% of the 2nd phase.

22. the high strength cold rolled steel plate according to claim 21 record is characterized in that: on the basis of aforementioned component, also contain the composition 1 group or 2 group or more of the following d group～g that represents with quality % among organizing, wherein,

D group: among Cu, Ni, Cr, the Mo more than a kind or 2 kinds, the total amount is below 1.0%;

E group: among Nb, Ti, the V more than a kind or 2 kinds, the total amount is below 0.1%;

F group: the B below 0.0030%;

G group: among Ca, the REM a kind or 2 kinds, the total amount is 0.0010～0.010%.

23. one kind has the r value: the manufacture method of the non-ageing high strength cold rolled steel plate of high r value more than 1.2 and good strain-aged hardening characteristics and normal temperature, it is characterized in that: % represents with quality, to contain C:0.025～0.15%, below the Si:1.0%, below the Mn:2.0%, below the P:0.08%, below the S:0.02%, below the Al:0.02%, N:0.0050～0.0250%, and N/Al is that the plate slab of the composition more than 0.3 is heated to slab heating temperature: more than 1000 ℃, carry out the roughing base that laminates, implement successively this thin slab execution finish rolling is gone out the material temperature: the finish rolling more than 800 ℃, and in the coiling temperature: make the hot-rolled process of hot-rolled sheet thereby reel below 800 ℃, thereby this hot-rolled sheet is implemented pickling and the cold rolling cold rolling process of making cold-reduced sheet, to this cold-reduced sheet in annealing temperature: more than the recrystallization temperature～implement pack annealing below 800 ℃, then at annealing temperature: Ac ₁Transformation temperature～(Ac ₃Transformation temperature-20 ℃) carries out continuous annealing under, then with speed of cooling: 10～300 ℃ of/second cold-reduced sheet annealing operations that are cooled to the humidity province below 500 ℃.

24. manufacture method according to the high strength cold rolled steel plate of claim 23 record, it is characterized in that: continue with the cooling after the aforementioned continuous annealing, carry out the overaging processing of residence time more than 20 seconds in the humidity province that aforementioned refrigerative cooling stops below the temperature, more than 350 ℃.

25. the manufacture method of high strength cold rolled steel plates according to claim 23 or 24 records is characterized in that: on the basis of aforementioned component, also contain the composition 1 group or 2 group or more of the following d group～g that represents with quality % among organizing, wherein,

D group: among Cu, Ni, Cr, the Mo more than a kind or 2 kinds, the total amount is below 1.0%;

E group: among Nb, Ti, the V more than a kind or 2 kinds, the total amount is below 0.1%;

F group: the B below 0.0030%;

G group: Ca, REM a kind or 2 kinds, the total amount is 0.0010～0.010%.

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