JP4441417B2 - High-tensile cold-rolled steel sheet with excellent formability and weldability and method for producing the same - Google Patents

Description

  The present invention has a maximum tensile strength (TS) of 700 MPa or more, a yield ratio of 0.60 or less, a Ceq value of 0.25 or less, excellent stretch formability and stretch flangeability, and excellent spot weldability and arc weldability. Further, the present invention relates to a high-tensile cold-rolled steel sheet that has all excellent impact resistance characteristics and is particularly suitable for structural members, reinforcing members, and suspension members for automobiles.

  In recent years, the use ratio of high-strength steel plates called high tension has been increasing in order to reduce the weight of automobile members. However, since the cold press formability decreases as the strength of the steel plate increases, the development of a steel plate having both high tensile strength and high formability has been desired. A steel plate of TS: 700 MPa or more, which is mainly used for structural members and reinforcing members of automobiles, is required to have excellent stretch flange formability, bend formability, and stretch formability. In general, stretch flange formability and bend formability are related to local elongation performance, so steel sheets with high local elongation and total elongation are oriented, while stretch formability is correlated with the n value of the steel sheet. Therefore, a steel sheet having a low yield ratio (TS / YS ratio) has been directed.

  Furthermore, since automobile parts are welded by spot welding, arc welding, laser welding, etc., it has been required to have high joint strength after welding. However, as the strength of steel sheets increases, the content of C, Si, Mn, etc. increases, and the strength of the weld zone decreases accordingly, resulting in low yield without increasing the content of alloying elements as much as possible. It has been desired to increase the strength while satisfying the ratio.

  In response to such a problem, Japanese Patent Application Laid-Open No. 11-279691 (Patent Document 1) discloses a high-strength galvannealed steel sheet having good workability with a TS of 490 to 880 MPa and a method for producing the same, and Japanese Patent Laid-Open No. 2000-80440 (Patent Document 2) discloses a cold-rolled thin steel sheet having excellent stretch flange characteristics and spot weldability of 700 MPa or more and a method for producing the same, and JP-A-2001-152287 (Patent Document). 3) shows a high-strength cold-rolled steel sheet excellent in spot weldability and a manufacturing method thereof, and JP-A-2001-226742 (Patent Document 4) shows a hot-dip galvanized steel sheet excellent in formability and a manufacturing method thereof. However, Japanese Patent Application Laid-Open No. 2003-231941 (Patent Document 5) discloses a high-strength cold-rolled steel sheet and a high-strength surface-treated steel sheet with good weldability and their manufacture Method is disclosed.

  However, in the method disclosed in Patent Document 1, if a TS of 700 MPa or more is to be obtained, the yield ratio is high, so that the stretchability is inferior, or if the yield ratio is to be lowered, C, Si, Mn is contained. It was necessary to increase the amount, and there was a problem that weldability deteriorated. Patent Document 2 discloses a cold-rolled thin steel sheet of TS: 700 MPa or more, which has excellent stretch flange characteristics and spot weldability, and further has strain age hardening, and a method for producing the same. However, in this method, since the main phase is bainite, the yield ratio is high, and since the strength-elongation balance is low, the bending formability is not always sufficient, and when trying to lower the yield ratio, C, Si, and Mn are contained. It was necessary to increase the amount, and as a result, there was a problem that weldability deteriorated.

  Further, Patent Document 3 discloses a high-strength cold-rolled steel sheet excellent in spot weldability with prescribed components and structure fractions and a method for producing the same. However, with this method, although the strength-elongation balance is very excellent, it is difficult to produce a steel plate of TS: 700 MPa or more with excellent weldability. Patent Document 4 discloses a hot-dip galvanized steel sheet having excellent formability with prescribed components and structure and a method for producing the same. However, although this method has a very good balance of strength and elongation, it is necessary to increase the amount of Si added to obtain TS: 700 MPa or more, and there is a problem that the chemical conversion processability deteriorates or the weldability deteriorates. Furthermore, there was a problem that the yield ratio was also high.

  Patent Document 5 discloses a high-strength cold-rolled steel sheet and a high-strength surface-treated steel sheet that have good weldability and require a TS of 780 MPa or more that defines the components and structure, and a method for producing the same. However, in this method, the ductility is low when the Si content is less than 0.4%, whereas the weldability deteriorates although the total elongation is slightly improved when the Si content is 0.4% or more. Had.

Japanese Patent Application Laid-Open No. 11-296991 JP 2000-80440 A JP 2001-152287 A JP 2001-226742 A JP 2003-231941 A

  The present invention has a maximum tensile strength (TS) of 700 MPa or more, a yield ratio of 0.60 or less, a Ceq value of 0.25 or less, excellent stretch formability, bending formability, stretch flangeability, and excellent spot weldability. -An object is to provide a high-tensile cold-rolled steel sheet suitable for automobile structural members, reinforcing members, and suspension members that have both arc weldability and excellent impact resistance.

  As a result of earnestly experimenting and studying to achieve the above object, the present inventors have added Ti and B, Mo or W in combination, and added other alloy elements in appropriate amounts, and ferrite. It has been found that a steel sheet having a tensile strength of 700 MPa or more and having a low yield ratio, high ductility, and excellent weldability can be produced by using a phase as a main component and the balance being martensite, bainite, and austenite. .

That is, the present invention is a high-tensile cold-rolled steel sheet excellent in formability and weldability, and the gist thereof is as follows.
(1) By mass%, C: 0.05 to 0.12%, Si: 0.4 to 1.5%, Mn: 1.0 to 3.0%, P: 0.04% or less, S: 0.01% or less, Ti: 0.003 to 0.05%, Nb: 0.04% or less, N: 0.01% or less, B: 0.0003 to 0.003%, and Mo : 1.0% or less, W: 0.03 to 1.0% in total of one or two of 1.0% or less, with the balance consisting of Fe and unavoidable impurities, the fraction of ferrite phase structure 60% to 95%, tensile strength: 700 MPa or more, yield ratio: 0.60 or less, and Ceq: 0.25 or less. High tensile cold rolling excellent in formability and weldability steel sheet.

(2) The high-tensile cold-rolled steel sheet having excellent formability and weldability according to (1), which includes one group or two or more groups of the following a group to c group in addition to the above components.
Group a: 0.002 to 0.1% by mass in total of one or more of V, Ta, and Al.
Group b: 0.001 to 0.01% by mass in total of one or more of Ca, Mg, Zr, Ce, and REM.
c group: 0.002-2.0 mass% in total of 1 type, or 2 or more types among Cr, Cu, and Ni.

(3) The precipitate particle size is 0.1 μm or more, and the density distribution of the iron carbide or borohydride having an fcc structure is 50,000 or less per 1 mm ² (1) or (2 ) High-tensile cold-rolled steel sheet with excellent formability and weldability as described above.
(4) Formability and weldability characterized by electroplating, hot dip galvanizing, or alloying galvanizing being applied to the steel sheet according to any one of (1) to (3) above. High-tensile cold-rolled steel sheet with excellent resistance.

(5) A steel material composed of the chemical component described in (1) or (2) above is heated to 1100 ° C. or higher, finish hot rolled at an Ar ₃ temperature or higher, and wound at 650 ° C. or lower. After cold rolling the hot-rolled steel sheet, in a continuous annealing facility, annealing is performed at an Ac ₁ temperature or higher and (Ac ₃ temperature −30 ° C.) or lower, and an average cooling rate between 650 ° C. and 450 ° C. is 10 to 200 ° C. A method for producing a high-tensile cold-rolled steel sheet excellent in formability and weldability, characterized in that it is cooled at / s and held in a temperature range between 400 ° C. and 250 ° C. for 100 to 400 s.

(6) The steel material comprising the chemical component described in (1) or (2) is heated to 1100 ° C. or higher, finish hot rolled at an Ar ₃ temperature or higher, and wound at 650 ° C. or lower. After cold rolling the hot-rolled steel sheet, it is annealed at a temperature not lower than Ac ₁ and not higher than (Ac ₃ temperature−30 ° C.) in a continuous hot dip galvanizing facility, and an average cooling rate between 10 ° C. and 450 ° C. is 10 to 10 ° C. A method for producing a high-tensile cold-rolled steel sheet having excellent formability and weldability, characterized by cooling at 200 ° C./s.

  The present invention provides a steel sheet that is suitable for structural members, reinforcing members, and suspension members for automobiles, and has high strength of 700 MPa or more with TS, excellent formability, and excellent weldability at low cost. it can. Furthermore, the steel sheet of the present invention is characterized by having both high bake hardenability and room temperature non-aging properties.

  First, the present inventors investigated the cause of failure to achieve TS: 700 MPa or more when DP alloy having a low yield ratio has a small additive amount of alloy elements such as C, Si, Mn, etc., which degrade weldability. did. As a result, the martensitic transformation or bainite transformation is completed during cooling after annealing of the continuous annealing line or continuous plating line, and the hard martensite phase or bainite phase is baked during aging in the overaging zone or alloying furnace. It has been found that TS: 700 MPa or more is difficult to obtain because of the back softening.

  As a result of studying a component system in which the temper softening of the hard phase does not occur, the present inventors have added one or two of Mo or W in addition to Ti and B, and set the other components in an appropriate range. As a result, it was found that the strength does not decrease even at a high overaging temperature or alloying temperature. Further, it was found that this component can achieve both excellent elongation characteristics.

  Next, the present inventors investigated the cause of excellent ductility. As a result, in steel containing Ti or B and further containing one or two of Mo or W, during annealing heating to the α + γ2 phase region or cooling after heating, on the old γ grain boundary or on the α / γ interface Precipitation of precipitated fcc structured coarse carbides or carbon borides has been remarkably suppressed, and this has been found to be one of the causes of excellent elongation properties, leading to the present invention.

The present invention is described in detail below.
First, the reasons for limiting the components will be described. In addition,% means the mass%.
C: C is mainly used for strength and structure adjustment. However, if it is less than 0.05%, it becomes difficult to achieve both a tensile strength of 700 MPa or more and moldability. On the other hand, if it exceeds 0.12%, it becomes difficult to ensure spot weldability. For this reason, the range was limited to 0.05 to 0.12%.

Si: Like C, Si: Si is mainly used for adjusting strength and structure. However, if it is less than 0.4%, the reduction in forming processability due to the deterioration of elongation becomes significant, and if it exceeds 1.5%, the wettability of plating decreases in the case of chemical conversion treatment or galvanized steel sheet. . Therefore, the Si content is limited to the range of 0.4 to 1.5%.
Mn: Similar to C, Mn is mainly used for adjustment of strength and structure. Since Mn is an element that enhances the hardenability of steel at a low cost, it is preferably used positively. However, if it is less than 1.0%, it becomes difficult to obtain a tensile strength of 700 MPa or more, and if it exceeds 3.0%, weldability deteriorates. For this reason, the appropriate range of Mn content was limited to the range of 1.0 to 3.0%. In addition, from the viewpoint of moldability, the content is preferably 2.4% or less.

  P: P tends to segregate in the central part of the plate thickness of the steel sheet and embrittles the weld. When the content exceeds 0.04%, the weld becomes brittle, so the appropriate range is limited to 0.04% or less. Although a minimum is not specifically limited, 0.0003% or more shall be contained as an unavoidable impurity. S: S exists as an impurity. Although it is mainly present in steel as MnS, when the S content exceeds 0.01%, the decrease in elongation becomes significant, so the range was limited to 0.01% or less. Although a minimum is not specifically limited, 0.0003% or more shall be contained as an unavoidable impurity.

  Ti: Ti is an important element in the present invention for fixing N in steel mainly as TiN, preventing precipitation of BN, and effectively expressing the combined effect of Mo or W and B as grain boundary segregation B. . However, if it is less than 0.003%, N is not sufficiently fixed and the effect of adding B is lost. Moreover, since elongation will fall by precipitation of TiC at 0.05% or more, the range was limited to 0.003-0.05%. A more preferable Ti content is in a range satisfying 3.0> ([% Ti] / [% N]) × (14/48)> 0.8, depending on the amount of N contained.

  Nb: Nb is used to increase the strength of the steel. However, when it exceeds 0.04%, the deterioration of elongation due to Nb carbide becomes remarkable. Therefore, the range is limited to 0.04% or less. From the viewpoint of manufacturing cost, it is more desirable to add Nb in an amount of 0.005% or more, and from the viewpoint of suppressing a decrease in elongation due to non-recrystallized ferrite, addition of 0.03% or less is desirable.

N: N is mainly used for controlling the crystal grain size in the austenite region. However, if N exceeds 0.01%, the amount of nitride deposited becomes excessive and the moldability is lowered, so the range of N content is set to 0.01% or less.
B: B is one of the most important elements in the present invention. However, if it is less than 0.0003%, it is difficult to obtain TS: 700 MPa or more, and if it exceeds 0.003%, formability deteriorates due to precipitation of coarse boride or boron carbide. For this reason, the appropriate range of B content was limited to the range of 0.0003 to 0.003%. Since B is an element that increases the hardenability by adding a small amount, it is desirable to add 0.0006% or more in terms of cost. Addition of 0.002% or less is more desirable from the viewpoint of suppressing material non-uniformity caused by macrosegregation of B.

Mo: Mo is one of the most important elements in the present invention, and has an effect of suppressing γ → α transformation after annealing and further suppressing temper softening of the hard phase. However, since elongation falls when it exceeds 1.0%, it limited to the range of 1.0% or less.
W: W is one of the most important elements in the present invention, and has the effect of suppressing the γ → α transformation after annealing and further suppressing the temper softening of the hard phase. However, since elongation falls when it exceeds 1.0%, it limited to the range of 1.0% or less.

  Total amount of Mo and W: If the total amount of addition of Mo and W is less than 0.03%, there is no effect of suppressing the γ → α transformation after annealing, and further precipitated during annealing or cooling after annealing. It is impossible to suppress the formation of fcc structure iron carbide or borohydride which deteriorates elongation. On the other hand, if the total exceeds 1.0%, the elongation decreases. For this reason, the appropriate range was limited to 0.03 to 1.0% in total of one or two of Mo and W. However, since the addition of a large amount of Mo and W increases the cost, it is desirable to contain the minimum amount necessary. In this case, the addition of 0.35% or less of the above total amount is a more preferable condition from the viewpoint that formation of fcc structure iron carbide or iron boride can be suppressed to the minimum necessary.

In the present invention, in addition to the above-described components, the object of the present invention can be achieved even when one group or two or more groups of group a to group c are contained.
Group a: 0.002 to 0.1% in total of one or more of V, Ta, and Al.
V, Ta, and Al are used as deoxidizing elements or carbonitride forming elements to adjust strength and structure. In order to acquire such an effect, it is preferable to contain 0.002% or more of 1 type or 2 types or more in total. However, if the total content exceeds 0.1%, the moldability deteriorates, so the total amount range was set to 0.002 to 0.1%.

b group: One or two of Ca, Mg, Zr, Ce, and REM in total 0.001 to 0.01%.
Ca, Mg, Zr, Ce, and REM are elements used for deoxidation, and it is preferable to contain one or two kinds in total in an amount of 0.001% or more. However, if the total content exceeds 0.01%, the moldability is deteriorated. Therefore, the total amount range is set to 0.001 to 0.01%. In the present invention, REM refers to La and lanthanoid series elements.

c group: 0.002-2.0 mass% of 1 type or the total of 2 or more types among Cr, Cu, and Ni.
Cr, Cu, and Ni are mainly used to adjust the strength of the steel sheet, and the effect is seen at 0.002% or more. However, when it exceeds 2.0%, surface property deterioration and molding processability are deteriorated, so the appropriate range of the total amount is limited to 0.002 to 2.0% or less. In addition, O (oxygen) may be included as 0.01% or less as another inevitable impurity.

  Structure fraction of ferrite phase: When the structure fraction of the ferrite phase is less than 60% or more than 95%, a yield ratio of 0.60 or less and excellent elongation characteristics are not compatible. For this reason, the structure fraction of the ferrite phase was limited to a range of 60 to 95%. The form of the ferrite phase includes recovered non-recrystallized ferrite and bainitic ferrite in addition to polygonal ferrite. The structure fraction of the ferrite phase is defined as an area ratio measured by image processing after observing the structure with a scanning electron microscope (SEM).

  The balance other than the ferrite phase is one or more of martensite, lath bainite, and austenite. Since the excellent elongation is partly due to the presence of the austenite phase, it is more preferable to contain 3% or more of the austenite phase. However, if it exceeds 9%, it becomes difficult to obtain a yield ratio of 0.60 or less, so the austenite phase fraction is preferably in the range of 3 to 9%. The recovered unrecrystallized ferrite and the lath bainite phase are separated by the crystal orientation mapping of the FESEM-EBSP method, and the ferrite phase has a continuously changing orientation within the grain or has a recovered fine cell structure. It is determined that the ferrite phase that has been used is an unrecrystallized ferrite phase. A method for quantifying the austenite fraction by X-ray diffraction is simple and suitable.

  In addition, the separation of the lath bainite phase, the martensite phase, and the tempered martensite phase is performed by microstructure observation with an SEM. Those in which no carbides exist in the lath are martensite phase, carbides are lined up in one direction in the lath, or those in which carbides are present between the laths are the lath bainite phase, and the parent phase Fe is contained in the lath. A carbide precipitated in the orientation variant is defined as tempered martensite. The crystal grain size of ferrite is not particularly limited, but it is preferably 7 μm or less in terms of nominal grain size from the viewpoint of balance of strength elongation.

  Yield ratio 0.60 or less: Yield ratio is a value defined by the ratio YS / TS of maximum tensile strength (TS) and yield strength (YS), and is an amount inversely proportional to the n value of the steel sheet. When the yield ratio is 0.60 or less, a good n value is obtained, and as a result, excellent stretch formability is realized. Therefore, the yield ratio is set to 0.60 or less in the present invention. In addition, it is more preferable that it is 0.55 or less as a necessary and sufficient value for the forming of many high-tensile steel plates.

Ceq value of 0.25 or less: In the present invention, weldability is evaluated by carbon equivalent Ceq in spot welding, which is most often used as a weldability index for steel sheets for automobiles. When Ceq exceeds 0.25, the joint tensile strength is lowered and the frequency of weld fracture at the time of forming is increased. Therefore, in the present invention, a steel plate having Ceq of 0.25 or less is provided as a steel plate having excellent weldability. Shall. More preferably, Ceq: 0.24 or less. The lower limit is not particularly limited, but is preferably 0.20 or more for securing TS.
Here, Ceq is a value defined by the following equation.
Ceq = C + (Si / 30) + (Mn / 20) + 2P + 4S (mass%)

Volume fraction of fcc (face-centered cubic) iron carbide or ferrocarbon borate: fcc structure iron carbide or borohydride precipitates at the old γ grain boundary or α / γ interface during annealing or cooling after annealing However, it is thought that local growth is reduced. When the precipitate diameter is less than 0.1 μm, there is almost no influence on the elongation characteristics, so the precipitate diameter was limited to 0.1 μm or more. Further, when the area fraction exceeds 50,000 per 1 mm ² , the local elongation characteristics deteriorate, so the appropriate range was limited to 50000 or less per 1 mm ² . 10,000 / mm ² or less is a more preferable range.

As a method for determining the distribution of the fcc-structured iron carbide or ferrocarbon borate, a method of extracting precipitates by the electrolytic extraction replica method and measuring the number per unit area by a transmission electron microscope is simple and suitable. In addition to fcc-structured iron carbide or iron-boride boride, there are iron carbides commonly called rhombohedral structure cementite in the steel sheet, but cementite has a different precipitation site from fcc-structured precipitates. It is known that the elongation is hardly deteriorated within the range of components within the range of, and cementite is excluded from the measurement.
In the present invention, a steel sheet having excellent total elongation and local elongation and excellent formability is provided. Specifically, TS (MPa) × total elongation (%) is 17000 or more, TS (MPa) × local. The thing which has the characteristic of 6000 or more by elongation (%) is pointed out.

Next, the reasons for limiting the production conditions of the steel sheet of the present invention will be described.
The slab used for hot rolling is not particularly limited. That is, what was manufactured with the continuous casting slab, the thin slab caster, etc. should just be used. It is also compatible with processes such as continuous casting-direct rolling (CC-DR) in which hot rolling is performed immediately after casting. The hot-rolled slab heating temperature needs to be 1100 ° C. or higher because it is necessary to redissolve carbonitride precipitated during casting or rough rolling.

When the finish rolling temperature is two phase region of the gamma + alpha, tissue heterogeneity in the steel sheet and the anisotropy of the material is increased, moldability after annealing so degraded, it was limited to the above Ar ₃ temperature.
The Ar ₃ temperature is calculated by the following formula.
Ar ₃ = 901-325 × C + 33 × Si-92 × (Mn + Ni / 2 + Cr / 2 + Cu / 2 + Mo / 2)

  The coiling temperature needs to be 650 ° C. or lower. If it exceeds 650 ° C., the non-uniformity of the structure after annealing becomes large and the moldability is lowered. From the viewpoint of making the microstructure after annealing fine and improving the strength ductility balance, it is more preferable to wind up at 600 ° C. or lower. The lower limit is not particularly required and may be room temperature, but is preferably 400 ° C. or higher from the viewpoint of reducing the cold rolling load. After winding, it is cooled, and then generally known treatments such as pickling are performed, followed by cold rolling. With regard to the cold rolling conditions, the number of rolling passes and the rolling reduction need not be specifically defined, and may be in accordance with ordinary methods.

What is necessary is just to follow a conventional method about the heating rate in a continuous annealing process or a continuous plating process. As for the maximum annealing temperature, if it is less than the Ac ₁ transformation point, it does not become a multiphase steel, and TS: 700 MPa or more cannot be obtained. On the other hand, if it exceeds (Ac ₃ temperature-30 ° C.), it becomes difficult to obtain a ferrite structure fraction of 60% or more, and as a result, the formability deteriorates and the yield ratio increases. Therefore the range of the annealing temperature Ac ₁ temperature or higher, and is limited to (Ac ₃ temperature -30 ° C.) within the range. From the viewpoint of increasing the ferrite fraction to 70% or more and further improving the moldability, the range of (Ac ₃ temperature −40 ° C.) or less is a more preferable condition.
The Ac ₁ temperature and Ac ₃ temperature can be estimated by the following equations.
Ac ₁ = 723 + 29Si-11Mn-17Ni + 17Cr + 7W (mass%)
^{_{Ac 3 = 910-203C 1/2 -30Mn + 45Si}} + 32Mo + 13W + 700P + 400Al + 400Ti-15Ni ( wt%)

  After the completion of the annealing, a process of rapidly cooling between 650 ° C. and 450 ° C. to form a hard phase is performed. If the average cooling rate during this period is less than 10 ° C./s, a hard phase with sufficient hardness cannot be obtained, and as a result, it becomes difficult to achieve both TS: 700 MPa or more and a yield ratio of 0.60. On the other hand, when the temperature exceeds 200 ° C./s, the material variation within the plate width increases, and the moldability may deteriorate. For this reason, the range of the average cooling rate was limited to the range of 10 to 200 ° C./s. From the viewpoint of further reducing the amount of additive alloy element from the viewpoint of weldability, it is more preferable to cool at a cooling rate of 20 ° C./s or more. As for the cooling method between 650 ° C. and 450 ° C., roll cooling, air cooling, water cooling, or any combination of these methods may be used.

When the cooling start temperature that limits the cooling rate is less than 650 ° C., sufficient hard phase hardness cannot be obtained, and it becomes difficult to achieve both TS: 700 MPa or more and a yield ratio of 0.60 or less. The upper limit of the starting temperature that limits the cooling rate is not particularly defined, but if it exceeds 750 ° C., the ferrite phase fraction tends to decrease and the elongation tends to decrease, so 750 ° C. or less is a preferable condition.
Regarding the lower limit of cooling that limits the cooling rate, if it exceeds 450 ° C., sufficient hard phase (martensite + bainite) hardness may not be obtained, and it is difficult to achieve both TS: 700 MPa or more and a yield ratio of 0.60. become. Although there is no particular limitation on the lower limit of the cooling stop temperature, it is preferably 300 ° C. or higher from the viewpoint of suppressing deterioration in formability due to strength variation in the steel sheet.

  Further, in the case of a continuous annealing line, heat treatment using an overaging zone is performed, but when the average plate temperature in the overaging zone is less than 250 ° C., a decrease in elongation is observed, and when it exceeds 400 ° C., TS: It is difficult to achieve a balance of 700 MPa or more, a yield ratio of 0.60 or less, and Ceq: 0.25 or less. Therefore, the appropriate range was limited to the range of 250 to 400 ° C., and the holding time was limited to 100 to 400 s. The upper limit of the holding time in the overaging zone is preferably within 360 s when the average plate temperature is less than 320 ° C. and within 300 s when the average plate temperature is less than 320 ° C.

  When hot dip galvanizing is performed, similarly, annealing between 650 ° C. and 450 ° C. is rapidly performed subsequent to the annealing heat treatment to form a hard phase. If the average cooling rate during this period is less than 10 ° C./s, a hard phase having sufficient hardness cannot be obtained, and as a result, it becomes difficult to achieve both TS: 700 MPa or more and a yield ratio of 0.60. On the other hand, when the temperature exceeds 200 ° C./s, the material variation within the plate width increases, and the moldability may deteriorate. For this reason, the range of the average cooling rate was limited to the range of 10 to 200 ° C./s. From the viewpoint of further reducing the amount of additive alloy element from the viewpoint of weldability, it is more preferable to cool at a cooling rate of 20 ° C./s or more.

  The plating bath temperature, plating component, alloying temperature, and time are not particularly limited, and may be performed by a generally known method. In addition, even if this cold-rolled steel sheet is electroplated, the material, formability, and weldability of the steel sheet are not impaired at all. That is, the steel sheet of the present invention is also suitable as a material for electroplating.

Next, the present invention will be described in detail with reference to examples.
A slab having the components shown in Table 1 was heated to 1200 ° C., subjected to hot rolling at a finish hot rolling temperature of 900 ° C., water-cooled in a water-cooled zone, and then wound up at a temperature shown in Table 2. . After pickling the hot-rolled sheet, it was cold-rolled at a cold reduction rate of 50 to 70% to obtain a cold-rolled sheet. After that, these cold-rolled sheets were subjected to annealing heat treatment under the conditions shown in Table 2, and after annealing, between 650 ° C. and 450 ° C. were cooled under the conditions of Table 2 and cooled to the overaging heat treatment temperature. Heat treatment was performed for 240 s under overaging conditions, followed by water cooling to room temperature. Finally, skin pass rolling was performed on the obtained steel sheet at a rolling reduction of 0.3%.

  For the steel sheet to be hot dip galvanized, after annealing, it is cooled between 650 ° C. and 450 ° C. under the conditions shown in Table 2, and then is passed through a galvanizing bath and brought to room temperature at a cooling rate of 10 ° C./s. Cooled down. For the alloying treatment, after passing through a galvanizing bath, the alloying treatment was performed at 500 ° C. for 30 s, cooled to room temperature at a cooling rate of 10 ° C./s, and finally obtained. The steel plate was subjected to skin pass rolling at a rolling reduction of 0.3%.

The obtained cold-rolled annealed plate or galvanized plate was subjected to a tensile test, microstructure and precipitate observation.
In the tensile test, a JIS No. 5 test piece was collected from a 1.2 mm thick plate (TD direction), and its YS, TS, total elongation, and local elongation were measured. The yield strength was measured by the 0.2% offset method. The bake hardenability test was conducted under the conditions of 2% pre-strain and 170 ° C. × 20 minutes. Moreover, the test which evaluates normal temperature non-aging property was also implemented by holding | maintaining a steel plate at 100 degreeC for 60 minutes, and performing a tensile test. All the BH contents of the inventive steel were 50 MPa or more, and the yield point elongation after holding at 100 ° C. for 60 minutes was 0.1% or less.

  The ferrite fraction was measured by SEM observation after corroding the structure with nital. When the structure was delicately determined, the structure was determined by analyzing the crystal orientation by the FESEM-EBSP method. Moreover, about the austenite fraction, it carried out by the X ray diffraction method. The distribution of fcc structure carbide or carbon boride was measured by observing the precipitate extracted in the carbon film by the electrolytic extraction replica method with a transmission electron microscope. Among these, the area distribution density of carbonized (boride) having an fcc structure and a particle diameter of 0.1 μm was measured by an electron diffraction method. The results of each test are shown in Table 2.

The present invention provides a steel sheet that is suitable for structural members, reinforcing members, and suspension members for automobiles, and has high strength of 700 MPa or more, excellent forming workability, and excellent weldability at a low cost. It can be expected to contribute greatly to the weight reduction of automobiles, and the industrial effect is extremely high.


Patent applicant: Nippon Steel Corporation
Attorney Attorney Shiina and others 1

Claims (6)

% By mass
C: 0.05 to 0.12%,
Si: 0.4 to 1.5%,
Mn: 1.0 to 3.0%
P: 0.04% or less,
S: 0.01% or less,
Ti: 0.003 to 0.05%,
Nb: 0.04% or less,
N: 0.01% or less,
B: 0.0003 to 0.003%
In addition,
Mo: 1.0% or less,
W: 0.03 to 1.0% in total of one or two of 1.0% or less, the balance is made of Fe and unavoidable impurities, and the structure fraction of the ferrite phase is 60 to 95% A high-strength cold-rolled steel sheet having excellent formability and weldability, characterized by tensile strength: 700 MPa or more, yield ratio: 0.60 or less, and Ceq: 0.25 or less. The high-tensile cold-rolled steel sheet having excellent formability and weldability according to claim 1, wherein, in addition to the components, one group or two or more groups of the following groups a to c are included.
Group a: 0.002 to 0.1% by mass in total of one or more of V, Ta, and Al.
Group b: 0.001 to 0.01% by mass in total of one or more of Ca, Mg, Zr, Ce, and REM.
c group: 0.002-2.0 mass% in total of 1 type, or 2 or more types among Cr, Cu, and Ni. 3. The forming process according to claim 1, wherein the particle diameter of the precipitate is 0.1 μm or more and the density distribution of the iron carbide or borohydride having an fcc structure is 50,000 or less per 1 mm ^2. High-tensile cold-rolled steel sheet with excellent weldability and weldability. The steel plate according to any one of claims 1 to 3, wherein the steel plate is electroplated, hot dip galvanized, or alloyed galvanized. steel sheet. The steel material comprising the chemical component according to claim 1 or 2 is heated to 1100 ° C or higher, finish hot rolled at an Ar ₃ temperature or higher, wound at 650 ° C or lower, and then cold rolled on the hot rolled steel sheet. After rolling, in a continuous annealing facility, annealing is performed at an Ac ₁ temperature or higher and an (Ac ₃ temperature −30 ° C.) or lower, and a temperature between 650 ° C. and 450 ° C. is cooled at an average cooling rate of 10 to 200 ° C./s. A method for producing a high-tensile cold-rolled steel sheet, which is excellent in formability and weldability, characterized by holding for 100 to 400 s in a temperature range between 400 ° C and 250 ° C. The steel material comprising the chemical component according to claim 1 or 2 is heated to 1100 ° C or higher, finish hot rolled at an Ar ₃ temperature or higher, wound at 650 ° C or lower, and then cold rolled on the hot rolled steel sheet. After rolling, in a continuous hot dip galvanizing facility, annealing is performed at a temperature not lower than Ac ₁ and not higher than (Ac ₃ temperature −30 ° C.), and cooled between 650 ° C. and 450 ° C. at an average cooling rate of 10 to 200 ° C./s. A method for producing a high-tensile cold-rolled steel sheet having excellent formability and weldability.