JP2007119899A - 490 MPa class low yield ratio cold-formed steel pipe excellent in weldability and manufacturing method thereof - Google Patents

Abstract

A cold-formed steel pipe having a tensile strength of 490 MPa or more and a low yield ratio without performing SR treatment, and a useful method for producing such a cold-formed steel pipe.
The steel sheet has a predetermined chemical composition, and the microstructure of the steel sheet is 4 to 70 area% of a polygonal ferrite phase, 0 to 20 area% of a pseudopolygonal ferrite phase, and 0 to 5 area%. When the aspect ratio (major axis / minor axis) is composed of a martensite phase of 4.0 or less and the balance is composed of a bainite phase, the thickness is t (mm) and the outer cold bending diameter is d (mm). It has a cold-formed part whose d is 10% or less.
[Selection figure] None

Description

  The present invention relates to a cold-formed steel pipe excellent in weldability and having a low yield ratio and a tensile strength of 490 MPa, and a method for producing the same, and particularly a building having a CFT (Concrete-Filled Tube) structure excellent in earthquake resistance. The present invention relates to a 490 MPa grade cold-formed steel pipe that can be suitably used for the present invention, and a useful method for producing such a cold-formed steel pipe.

  Building structures are required to have excellent seismic resistance and fire resistance. To build buildings with the above-mentioned CFT structure, which are particularly excellent in earthquake resistance, they exhibit excellent weldability with high strength and low yield ratio. Cold forming steel pipe is required.

  Various techniques for cold forming steel pipes that satisfy these required characteristics have been proposed so far. For example, in Patent Document 1, steel plates that have been subjected to hot rolling and then air-cooled or water-cooled for 600 MPa class and 800 MPa class low yield ratio steel pipes are expressed as t / D (t: plate thickness, D: steel pipe outer diameter) ≦ Technology to manufacture steel pipes by cold forming in the range of 10% and normalize the steel sheet controlled to the yield ratio (YR) ≦ 80−0.8 × t / D in the temperature range of 750 to 850 ° C. Is disclosed.

Further, in Patent Document 2, a steel pipe having a low yield ratio of 590 MPa class is subjected to rolling so that the rolling finish temperature is (Ar ₃ −20 ° C.) to (Ar ₃ + 120 ° C.), and then the steel plate ( Ar ₃ −100 ° C.) to (Ar ₃ −120 ° C.), followed by quenching immediately from this temperature to room temperature, and further tempering to a temperature range of Ac ₁ point or less, and the above t / D ≦ 10% Is cold formed into a tube in the range of, and then annealed in a temperature range of 500 to 600 ° C.

Furthermore, in Patent Document 3, for a 590 MPa class low yield ratio steel pipe, it is reheated to a temperature of Ac ₃ point or higher and quenched or tempered and tempered, and cooled in the above range of t / D ≦ 10%. It is disclosed that a steel pipe is manufactured by performing inter-forming and then reheated to a temperature range of 650 to 750 ° C. to normalize.

  Each of the above technologies is directed to a cold-formed steel pipe having a low yield ratio of 590 MPa, and among them, Patent Document 1 performs normalization after cold forming. In patent document 2, since it cools to a two-phase area | region on a rolling line, it will cause the productivity fall in rolling and is not preferable from an economical viewpoint.

  In the technique of the above-mentioned patent document 3, since the steel material contains elements such as Cu and Ni as essential components, there is a problem that the material cost is increased. In this technique, the strength of the steel pipe is improved by precipitation strengthening due to the addition of Cu, but the temperature on the outer surface side and the inner surface side becomes uneven in the heat treatment process, causing unevenness in the precipitation of Cu. It is expected that it will occur.

In any of the above techniques, after cold forming, it is necessary to perform heat treatment for the purpose of reducing the yield ratio, and there are problems in terms of cost and productivity. In addition, when the so-called Delay DQ method of Patent Document 2 is applied, the strength as quenched is lower than direct quenching (DQ) from a temperature of _three or more points of Ar, so the alloy element is increased to compensate for it. As a result, weldability deteriorates.

Therefore, as a method of not performing heat treatment after cold forming, a technique such as Patent Document 4 has also been proposed. In this technique, after hot rolling, it is reheated to Ac _{3 to} 1000 ° C. and quenched. Using a steel plate that is reheated to 700 to 850 ° C. and quenched, tempered at Ac ₁ point or less, and controlled to YR (%) ≦ 80−0.8 × t / D, t / D A steel pipe is manufactured by cold forming in a range of ≦ 10%, thereby obtaining a low yield ratio 600 MPa class steel pipe for construction having a sheet thickness of 100 mm or less and a YR in the pipe axis direction of 80% or less. .

  This technology is intended for 600MPa class low-yield ratio steel pipes to prevent softening due to reheating and quenching, steel pipe toughness improvement, welding, stress relief treatment, etc. Is an essential process, and some problems remain in terms of productivity and cost. Moreover, this technique requires an increase in the amount of alloy elements from the viewpoint of securing strength, and still has a problem in terms of weldability.

On the other hand, for example, a technique as disclosed in Patent Document 5 has been proposed as a method for producing a 490 MPa class low yield ratio high tensile steel sheet. In this technique, rolling is performed so that the cumulative rolling reduction at 900 ° C. or less is 50% or more, and rolling is finished at Ar ₃ points or more, cooling to Ac ₁ point or less, and then within a range of 730 to 850 ° C. or less. It is reheated and air cooled.

In this technique, since the strength is lower than that obtained by quenching (Q ′) treatment from a two-phase region temperature (Ac ₁ point and less than Ac ₃ point), the carbon equivalent Ceq (JIS) is 0.00. Those with good weldability at 40% or less can be applied to plate thicknesses up to about 32 mm (for example, steel Nos. 1, 2, 4 to 6 in Table 1), but are applied to thick steel pipes for cold forming. If it is going to be, it is necessary to raise carbon equivalent Ceq significantly (for example, No. 3 of Table 1), weldability deteriorates in connection with it, and preheating is needed. Moreover, in order to increase the rolling reduction in the austenite non-recrystallized region (about 900 to Ar ₃ points), the requirement that the acoustic anisotropy required for the steel for construction is small is not satisfied.
Japanese Patent Application Laid-Open No. 6-128641 Japanese Patent No. 2529042 Patent Claim etc. JP, 7-233416, A Claims etc. Japanese Patent Laid-Open No. 7-109521 Japanese Patent Laid-Open No. 55-115921 Claims, Table 1, etc.

  By the way, with the revision of the new seismic design method (1981), the design concept of allowing the plastic deformation of steel materials in the case of a large earthquake and absorbing the energy of the earthquake to prevent the collapse of the structure in the building field As a result, a low yield ratio has been required as a necessary characteristic for steel materials.

  In cold-formed steel pipes applied to concrete-filled steel pipe columns, when severe cold bending of t / D: 5 to 10% is applied, a strain (ε) equivalent to about 2.5 to 5% at t / 4 part Therefore, even if it is a steel material having a tensile strength of 490 MPa class, the target yield ratio (yield point / tensile strength) of 85% or less cannot be ensured. In such a case, annealing (Stress Relieving: SR treatment) for the purpose of removing residual stress after cold forming is unavoidable, resulting in higher costs, longer construction periods, and lower productivity.

  The present invention has been made under such circumstances, and an object of the present invention is to provide a cold-formed steel pipe having a tensile strength of 490 MPa or more and a low yield ratio without performing SR treatment, and such a cold-formed steel pipe. It is to provide a useful method for manufacturing.

The 490 MPa class low yield ratio cold-formed steel pipe of the present invention capable of achieving the above object is C: 0.07 to 0.18% (meaning of mass%, the same shall apply hereinafter), Si: 0.05 to 1. 0%, Mn: 0.7 to 1.7% (provided that the ratio of Mn content [Mn] to C content [C] [Mn] / [C] ≦ 15), Ti: 0.002 to 0.8%. 025%, sol. In addition to containing Al: 0.005 to 0.1% and N: 0.001 to 0.008%, Cr: 0.6% or less (including 0%), Mo: 0.5% or less (0 %) And V: 0.08% or less (including 0%), including one or more selected from the group consisting of carbon equivalent Ceq values of 0.34 to It is within the range of 0.42%, and the A value shown by the following formula (2) satisfies 1.1 to 2.6, and the balance is made of a steel plate having a chemical composition composed of Fe and inevitable impurities. And the microstructure of the steel sheet is 4 to 70 area% of polygonal ferrite phase, 0 to 20 area% of pseudopolygonal ferrite phase, and 0 to 5 area%, and the aspect ratio (major axis / minor axis) is 4. 0.0 or less island martensite phase, the balance is composed of bainite phase, It has a gist in that it has a cold forming portion where t / d is 10% or less when t (mm) and the outer cold bending diameter is d (mm).
Ceq = [C] + [Si] / 24 + [Mn] / 6 + [Ni] / 40 + [Cr] / 5
+ [Mo] / 4 + [V]) / 15 (1)
However, [C], [Si], [Mn], [Ni], [Cr], [Mo] and [V] are the contents (mass of C, Si, Mn, Ni, Cr, Mo and V, respectively). %).
A = (2.16 {Cr} +1) × (3.0 {Mo} +1) × (1.75 {V} +1) (2)
However, {Cr}, {Mo} and {V} indicate the amount of solid solution (mass%) in the steel sheet in Cr, Mo and V, respectively.

  In the cold-formed steel pipe of the present invention, if necessary, (a) Cu: 0.5% or less (not including 0%) and / or Ni: 3.0% or less (not including 0%), ( b) Nb: 0.015% or less (not including 0%), (c) Ca: 0.003% or less (not including 0%), (d) Rare earth element: 0.02% or less (0% It is also effective to contain (not contained), etc., and the characteristics of the steel pipe can be further improved according to these contained components.

In producing the cold-formed steel pipe as described above, a steel piece having a chemical composition defined in the present invention is heated to a temperature range of 950 to 1250 ° C., and austenite-free steel represented by the following formula (3) is used. After the rolling was finished by setting the cumulative reduction ratio below the recrystallization temperature A _γ (° C.) to 60% or less (including 0%) to form a steel plate, the temperature from the Ar ₃ transformation point or higher to 450 ° C. or lower was 4 to 4 ° C. Accelerated cooling at a cooling rate of 100 ° C./second, then reheating to a temperature range of 730 to 830 ° C., followed by quenching, followed by cold forming with a t / d of 10% or less. good.
A _γ (° C.) = 887 + 467 [C] + (6445 [Nb] −644√ [Nb]) +
(732 [V] -230√ [V]) + 890 [Ti] +363 [Al] -357 [Si]
... (3)
However, [C], [Nb], [V], [Ti], [Al], and [Si] indicate the contents (mass%) of C, Nb, V, Ti, Al, and Si, respectively.

  In this manufacturing method, (1) after reheating to a temperature range of 730 to 830 ° C. and quenching, the steel sheet is tempered at 500 ° C. or less. (2) After the rolling is finished, accelerated cooling is performed. Prior to, it is preferable to add conditions such as performing online leveler correction, (3) cold forming with a steel plate temperature of 400 ° C. or lower.

  According to the present invention, by appropriately adjusting the chemical composition of the steel sheet and appropriately controlling the volume fraction of each phase in the microstructure, the 490 MPa class with a low yield ratio can be obtained without performing SR treatment. A cold-formed steel pipe can be obtained, and such a cold-formed steel pipe is obtained by appropriately controlling manufacturing conditions, and the obtained steel pipe can be suitably used for a building having a CFT structure.

  The present inventors have a cold-formed portion where t / d is 10% or less when the plate thickness is t (mm) and the outer cold bending diameter is d (mm), and the tensile strength is 490 MPa. In the above steel pipe, in order to achieve excellent weldability and to achieve a yield ratio of 85% or less, which is a target value, the chemical composition and the microstructure were examined in detail.

As a result, in order to reduce the yield ratio in the steel pipe, the yield ratio at the steel plate stage is lowered in advance by an amount higher than the increase in the steel pipe, and the uniform elongation δ _u (plastic elongation up to the maximum load) is increased. It was found that this is important.

In order to achieve both a low yield ratio and tensile strength at the steel plate stage, the microstructure is divided into a bainite phase (B) that becomes a hard phase and a polygonalized ferrite phase that becomes a soft phase (polygonal ferrite phase: α It was proved effective to coexist _{P 2} ) and control the area fraction of the polygonal ferrite phase (α _P ) to 40 to 70%. It was also found that the uniform elongation δ _u can be increased by polygonalizing the ferrite phase.

In the steel plate stage (before cold forming), when martensite is formed in an island shape, the yield point is reduced and the yield ratio after cold bending is further lowered. Island-like martensite consists of a mixture of martensite and austenite phases (residual austenite phase) (Martensite-Austenite constitute: MA phase), but the retained austenite phase present in island-like martensite. gamma _R can further increase the uniform elongation [delta] _u by transformation to deformation-induced martensite by cold bending. When such a structure is cold-bent to form a cold-formed steel pipe, the retained austenite phase in the structure disappears and exists as a transformed martensite phase.

  In the cold-formed steel pipe of the present invention, it is necessary to appropriately control the microstructure from the above viewpoint, and the reasons for limiting the range (area fraction) of each phase in this structure are as follows.

[Polygonal ferrite phase α _P : 40 to 70 area%]
In order to reduce the yield ratio, it is effective to generate polygonalized ferrite (α _P ) having a low dislocation density in the microstructure after transformation, and in order to lower the yield ratio in advance in the steel plate stage, It is necessary to control the area fraction within a range of 40 to 70%. When the area fraction of the polygonal ferrite phase (α _P ) exceeds 70%, it becomes difficult to secure the target strength in the thick material. On the other hand, when the area fraction of the polygonal ferrite phase (α _P ) is less than 40%, the yield ratio exceeds the target value (85%).

[Pseudopolygonal ferrite phase α _q : 0 to 20 area%]
The pseudopolygonal ferrite phase (α _q ) having a high dislocation density increases the strength, while preventing the movement of movable dislocations and increasing the yield ratio. Therefore, it is better as little as possible, and the area fraction is about 0 to 20%. There is a need to. Preferably, the content is about 0 to 15%.

[Islandic martensite phase MA: 0 to 5 area%]
The martensite phase (M) or martensite-austenite mixed phase (MA phase) in the steel plate stage is locally island-like without bainite transformation in the part where C and alloy elements are largely segregated in untransformed austenite. In the martensite phase (M) and the retained austenite phase (γ _R ). Of these, the martensite phase (M) effectively acts to increase the tensile strength and reduce the yield ratio. Residual austenite (γ _R ) effectively works to increase the uniform elongation δ _{u in} order to develop a processing-induced transformation due to external processing strain. Therefore, in cold-formed steel pipe, in order to further promote the reduction in yield ratio and the increase in uniform elongation δ _u , the island-like martensite phase (MA: martensite phase M after transformation from retained austenite M Including). The area fraction of the island martensite phase (MA) is preferably about 0 to 5%. If the area fraction of the island martensite phase (MA) exceeds 5%, the toughness will deteriorate. This area fraction is preferably about 0 to 4%.

[Aspect ratio of island martensite phase (MA): 4.0 or less]
Even if the area fraction of the island martensite phase (MA) is 5% or less, if the aspect ratio (major axis / minor axis) exceeds 4.0 in the shape, the uniform elongation δ _u increases. The toughness is also deteriorated. Further, since the MA phase is formed at the prior austenite grain boundaries, controlling the aspect ratio to 4.0 or less is a consequence of the low degree of expansion of the prior austenite grains, and the rolling texture Therefore, the acoustic anisotropy of the seam welded portion of the steel pipe (corresponding to a non-machined portion of the end bend) can be reduced.

  The cold-formed steel pipe of the present invention has a cold-formed portion where t / d is 10% or less when the plate thickness is t (mm) and the outer cold bend diameter is d (mm). In cold working such that t / d exceeds 10%, the yield ratio on the tensile deformation side exceeds 85% after working. Therefore, in order to suppress the increase in yield ratio, Or a stress-relieving annealing process (the SR process) after the molding is required. Therefore, the t / d needs to be 10% or less. This t / d is preferably 7.5% or less. The processing method for achieving this t / d is not limited to press bending, and for example, roller bending, compression pressing, spinning, and the like can be applied. The bending temperature is acceptable not only at room temperature but also up to a temperature that does not impair the material of the steel sheet of the present invention (about 400 ° C.). The cold-formed steel pipe of the present invention includes both circular and square cross-sectional shapes. The outer cold bend diameter d means a curvature diameter at a cold-formed (bent) portion. When the cross-sectional shape of the steel pipe is circular, the outer cold bend diameter d matches the outer diameter D of the steel pipe. Will do.

In the cold-formed steel pipe of the present invention, in order to control the quantitative ratio of ferrite (α _p ) in the microstructure as described above (area fraction: 40 to 70%), the ferrite nose of the transformation curve is set to the short time side. Specifically, the ratio of the Mn content [Mn] to the C content [C] ([Mn] / [C]) is set to 15 or less, and the temperature in the two-phase region (α + γ region) is maintained. It is effective to facilitate the two-phase separation of C. In order to exert such effects, it is effective to set the two-phase region quenching temperature to about 730 to 830 ° C. (this condition will be described later).

  The softening of ferrite and the hardening of cementite add a negative segregation element for ferrite and diffuse the negative segregation element for ferrite into untransformed austenite in the coexistence state of austenite and ferrite existing in two-phase temperature holding. Then, it was considered effective to cause bainite transformation and to concentrate the exhaled alloy elements in cementite during the bainite transformation process.

Focusing on the fact that Cr, Mo, and V have large effects as negative segregation elements for ferrite, the amount defined by equation (2) is controlled to 1.1 to 2.6 as the amount of these solid solutions. By doing so, the alloy element can be separated into two phases. Incidentally, the (2) in order to control the 1.1 to 2.6 the amount defined by the equation, and heating the billet to a temperature range of: 950 ° C., Ar ₃ transformation point after the end of rolling It is effective to quench Cr, Mo and V in a solid solution state and then quench in a two-phase region while avoiding precipitation in the nitride precipitation temperature region of each element by rapid cooling from the above temperature ( This condition will be described later).

For lowering the yield point and increasing the uniform elongation δ _u , it is effective to lower the dislocation density of the ferrite by grain growth without ferrite after the ferrite transformation and making it polygonal. is there.

  In order to reduce the yield ratio of the cold-formed steel pipe and to secure toughness, it is necessary to form a more isotropic martensite phase. From this viewpoint, the aspect ratio must be 4.0 or less. is there. In addition, since the rolling texture is reduced by reducing the aspect ratio, it is effective in reducing the acoustic anisotropy at the seam welded portion of the steel pipe (corresponding to the unmachined portion of the end bend).

As a means for suppressing the flattening of the martensite phase or the mixed layer with austenite at the steel plate stage, the cumulative rolling reduction at the austenite non-recrystallization temperature A _γ or less represented by the above formula (3) is 60% or less. Ending the rolling as described above is effective in suppressing flattening of the prior austenite grains and the martensite phase or grain mixed with the austenite. In addition, the rolling end temperature at this time is preferably _{equal to} or higher than the austenite non-recrystallization temperature Aγ from the viewpoint of suppressing flattening of the prior austenite grains. Furthermore, in order to control the martensite structure fraction as described above, it is necessary to replace C in the carbon equivalent formula with an alloy element (particularly, Cr, Mo, V, etc.) shown in the present invention.

In cold-formed steel pipe of the present invention, the microstructure, other than the above (balance), after it is made of essentially bainite, is obtained by precipitation of polygonal ferrite alpha _P of the present invention ranges Therefore, Immediate acceleration cooling may be performed so as not to cause pearlite transformation.

  By the way, in the cold-formed steel pipe of the present invention, it is necessary that the weldability is good. To that end, by adding B to the martensite in the weld heat affected zone (HAZ), or Bainite can be suppressed and crack resistance and HAZ toughness can be improved. Further, by adding Ti, TiN is generated, and the toughness is improved by exerting the effect of refining the prior austenite grains in the base material and HAZ.

  Next, the reason for limiting the chemical component composition in the cold-formed steel pipe of the present invention will be described. First, in the present invention, as described above, C: 0.07 to 0.18%, Si: 0.05 to 1.0%, Mn: 0.7 to 1.7% (provided that Mn content [Mn] And C content [C] ratio [Mn] / [C] ≦ 15), Ti: 0.002 to 0.025%, sol. Al: 0.005 to 0.1% and N: 0.001 to 0.008%, Cr: 0.6% or less (including 0%), Mo: 0.5% or less (0%) And V: 0.08% or less (including 0%), and one or more selected from the group consisting of 0 and 8%, and appropriate values defined by the above formulas (1) and (2) Although it is necessary to control within such a range, the reasons for limiting the ranges of these elements are as follows.

[C: 0.07 to 0.18%]
C is the cheapest element and is effective for increasing the strength. However, if excessively contained, the weldability is remarkably lowered, so the upper limit of the content is 0.18%. However, when the C content is less than 0.07%, the strength is insufficient, and in order to make up for it, addition of alloy elements is required. However, if these alloy elements are added excessively, the yield ratio is increased. Increase, which is not preferable. In order to secure the target strength (tensile strength of 490 MPa or more) while suppressing the increase in yield ratio, C needs to be contained at least 0.07% or more. In addition, from a viewpoint of coexistence of base material strength and welded HAZ toughness, a preferable lower limit of the C content is 0.08%, and a preferable upper limit is 0.16%.

[Si: 0.05-1.0%]
Si needs to be contained in an amount of 0.05% or more for deoxidation. However, if it exceeds 1.0%, the weldability and the HAZ toughness are deteriorated. For these reasons, the Si content needs to be 0.05 to 1.0%. In addition, the minimum with preferable Si content is 0.10%, and a preferable upper limit is 0.9%.

[Mn: 0.7 to 1.7% (provided that the ratio of Mn content [Mn] to C content [C] [Mn] / [C] ≦ 15)]
Mn is effective as an element that increases both strength and toughness. In order to exhibit such an effect, it is necessary to contain 0.7% or more of Mn. However, if Mn is contained excessively, the weldability and the HAZ toughness deteriorate significantly, so the upper limit is made 1.7%. In addition, the minimum with preferable Mn content is 1.0%, and a preferable upper limit is 1.6%.

  Further, the Mn content needs to be adjusted to an appropriate range in relation to the C content. The ratio [Mn] / [C] between the Mn content [Mn] and the C content [C] is the overhang of the ferrite transformation curve (nose) in the continuous cooling transformation curve (CCT curve) and isothermal transformation curve (TTT curve). ) Is a factor that controls the degree of component), and when the ratio [Mn] / [C] exceeds 15, the ferrite nose retreats to the long time side, so that the two-phase region heat treatment (Q ′) Thus, the holding time for obtaining an equilibrium two-phase structure (α + γ) becomes long, and it becomes inefficient due to restrictions on production. Therefore, the ratio [Mn] / [C] needs to be 15 or less.

[Ti: 0.002 to 0.025%]
Ti exists as fine TiN in the steel during slab heating, and has the effect of preventing coarsening of the heated austenite grains. Combined with proper austenite (γ) recrystallization temperature rolling, subsequent γ non-recrystallization temperature A _γ region rolling, and fine TiN formation, it is possible to ensure good toughness and ultrasonic acoustic anisotropy. It is. The Ti is a TiN as ferrite transformation nuclei from reverse transformed austenite in the Q 'treatment after direct quenching, by promoting the precipitation of polygonal ferrite, it is effective to yield ratio reduction, increase of uniform god beauty [delta] _u . In order to exert such effects, the Ti content needs to be 0.002% or more. However, even if Ti is contained excessively, the effect is saturated, so the upper limit is made 0.025%. In addition, the minimum with preferable Ti content is 0.008%, and a preferable upper limit is 0.015%.

[Sol. Al: 0.005 to 0.1%]
Al needs to be contained at least 0.005% for deoxidation, but if it is contained excessively, nonmetallic inclusions increase and toughness decreases, so it is necessary to make it 0.1% or less. . In addition, the minimum with preferable Al content is 0.01%, and a preferable upper limit is 0.06%.

[N: 0.001 to 0.008%]
N reacts with Ti to produce TiN, and is an element effective in preventing austenite coarsening during heating. In order to exert such an effect, it is necessary to contain at least 0.001% or more, but if contained excessively, the toughness of the welded joint portion deteriorates, so it is necessary to make it 0.008% or less. In addition, the minimum with preferable N content is 0.002%, and a preferable upper limit is 0.006%.

[Cr: 0.6% or less (including 0%), Mo: 0.5% or less (including 0%) and V: 0.08% or less (including 0%) Or two or more kinds [and the amount of solid solution satisfying the formula (2)]
Cr, Mo, and V are elements that improve the strength. However, when precipitated as a compound, the yield ratio is increased by precipitation strengthening, while the toughness is deteriorated. In order to ensure high strength and high toughness while keeping the yield ratio low, it is effective to cause segregation to be positive segregation to cementite and negative segregation to ferrite in a solid solution state. Therefore, the contents of Cr, Mo and V are 0.6% or less, 0.5% or less and 0.08% or less (all include 0%), respectively. It is necessary to control within the range of 1.1 to 2.6 with the A value defined by the equation. When the amount of component elements and the A value exceed the upper limit, weldability is hindered. On the other hand, if the A value is less than 1.1, the yield ratio after forming the steel pipe cannot satisfy the target value. Each element is preferably Cr: 0.3% or less, Mo: 0.3% or less, and V: 0.06% or less. Moreover, the preferable range of A value is about 1.05-2.0.

[Ceq: 0.34 to 0.42%]
The carbon equivalent Ceq represented by the formula (1) is an index representing the curability of the HAZ (for example, JIS G 3106), reduces the weld crack sensitivity, and reduces the crack prevention preheating temperature in the y-type weld crack test to 25. In order to set the temperature at or below ° C, the Ceq value needs to be set to 0.42% or less. On the other hand, in order to ensure a tensile strength of 490 MPa or more, the Ceq value needs to be 0.34% or more. The preferable lower limit of the Ceq value is 0.36%, and the preferable upper limit is 0.40%. In addition to the basic components C, Si, Mn, Cr, Mo, V, etc., the above formula (1) includes the component (Ni) contained as necessary in the formula. When the component is contained, it may be calculated as the value of the formula (1) in consideration of the content, and when not contained, it may be calculated without considering these contents.

  In the cold-formed steel pipe of the present invention, in addition to the above components, it is composed of Fe and unavoidable impurities, but may contain a trace component (acceptable component) inevitably mixed for melting (for example, P, S, O, B ≦ 0.0005%, etc.), such steel slabs are also included in the scope of the present invention. The cold-formed steel pipe of the present invention may further include (a) Cu: 0.5% or less (not including 0%) and / or Ni: 3.0% or less (not including 0%) as necessary. (B) Nb: 0.015% (not including 0%), (c) Ca: 0.003% or less (not including 0%), (d) Rare earth element: 0.02% or less (0%) It is also effective to contain these components, etc., but the reason for limiting the range when these components are contained is as follows.

[Cu: 0.5% or less (not including 0%) and / or Ni: 3.0% or less (not including 0%)]
Since these elements are expensive and increase the yield ratio, addition of these elements is preferably avoided as much as possible. However, since a thick steel plate has an effect of suppressing the strength reduction in the central portion of the plate thickness, a small amount may be added. When these elements are added, it is necessary to contain Cu up to 0.5% and Ni up to 3.0%. A more preferable upper limit of the Cu content is 0.3%, and a more preferable upper limit of Ni is 1.5%.

[Nb: 0.015% (excluding 0%)]
Nb is known as an element that improves both strength and toughness. However, when accelerated cooling is performed after hot rolling, the amount of bainite in the second phase structure is obtained in steel containing Nb, which is a hardenability improving element. Increases, and it becomes difficult to produce soft ferrite. As a result, the yield ratio increases. For these reasons, when Nb is contained, the content is preferably up to about 0.015%. The upper limit with more preferable Nb content is about 0.010%.

[Ca: 0.005% or less (excluding 0%)]
Ca has a spheroidizing effect of nonmetallic inclusions and is effective in reducing anisotropy. However, when Ca is contained in an amount exceeding 0.005%, the toughness deteriorates due to an increase in inclusions. Therefore, when Ca is contained, the content is preferably 0.005% or less. The preferable lower limit of the Ca content is 0.0005%, and the more preferable upper limit is 0.003%.

[Rare earth elements: 0.02% or less (excluding 0%)]
A rare earth element (hereinafter abbreviated as “REM”) is an element that suppresses abnormal austenite growth in the presence of TiN as oxysulfide and improves the toughness of HAZ, but exceeds 0.02% in excess. If done, the cleanliness of the steel is deteriorated and internal defects are generated. In order to exhibit the effect by REM, it is preferable to contain 0.002% or more, and a more preferable upper limit is 0.01%. As REM, any of scandium (Sc), yttrium (Y) and lanthanoid series rare earth elements belonging to Group 3 of the periodic table can be used.

  In order to produce the cold-formed steel pipe of the present invention, basically, using a steel piece from a slab produced by a continuous casting method or an ingot-making method, a process such as heating-hot rolling-cooling-heat treatment, Alternatively, a steel pipe having the above chemical composition and structure may be manufactured through a process such as controlled cooling after hot rolling (including accelerated cooling and direct quenching). Although it does not specifically limit (Experiment No. 42-46 of an after-mentioned Example), manufacturing according to this invention method is preferable. Next, each requirement prescribed | regulated with the manufacturing method of this invention is demonstrated.

[Heating temperature of steel slab: 950 to 1250 ° C.]
When the heating temperature of the steel slab exceeds 1250 ° C., the austenite grains of the steel slab undergo rapid grain growth, the transformed structure becomes a coarse bainite structure, and the toughness of the steel sheet becomes extremely low. On the other hand, when the heating temperature is less than 950 ° C., the cumulative reduction ratio below (γ non-recrystallization temperature Aγ−50 ° C.) increases, and excessive austenite grain refinement occurs, yield point YP, 0 The 2% yield strength σ _0.2 and the yield ratio YR will be significantly increased. For these reasons, the heating temperature of the steel slab needs to be in the range of 950 to 1250 ° C. The heating temperature is preferably 1000 ° C. or higher and 1150 ° C. or lower.

[Cumulative rolling reduction at γ non-recrystallization temperature A _γ or less is 60% or less]
As described above, in order to suppress the flattening of the mixed phase with the martensite phase or austenite at the steel plate stage, the cumulative rolling reduction at the γ non-recrystallization temperature A _γ or less needs to be 60% or less. . On the other hand, when the cumulative rolling reduction exceeds 60%, excessive austenite grain refinement occurs and the yield ratio increases. The “rolling ratio” is {(t ₁ −t ₂ ) / t ₁ } × 100 when the thicknesses of the steel sheet before and after rolling are t ₁ (mm) and t ₂ (mm), respectively. (%).

[After the rolling is completed, cooling is performed at a cooling rate of 4 to 100 ° C./second from the Ar ₃ transformation point to 450 ° C.]
For the purpose of ensuring uniform dispersion of C in the microstructure of the steel sheet and solid solution of Cr, Mo, V and ensuring the strength, it is necessary to accelerate cooling from the Ar ₃ transformation point to 450 ° C. after rolling. is there. If the cooling start temperature is lower than the Ar ₃ transformation point, the cooling stop temperature is higher than 450 ° C., or the cooling rate is less than 4 ° C./second, transformation strengthening is insufficient, Cr, The total solid solution of Mo and V is not achieved. The upper limit of the cooling rate at this time needs to be 100 ° C./second or less from the viewpoint of the limit of the cooling capacity of the cooling medium. In the present invention, the Ar ₃ transformation point is a value calculated by the following equation (4).
Ar ₃ transformation point = 910-310 [C] -80 [Mn] -20 [Cu] -15 [Cr]-
55 [Ni] -80 [Mo] +0.35 (t-8) (4)
Where t: thickness

[Quenching after reheating to 730-830 ° C temperature range]
By maintaining the accelerated cooled steel sheet at a two-phase region (α + γ) temperature, once dispersed C is two-phase separated into ferrite and reverse-transformed austenite, causing negative segregation of C in ferrite and positive segregation in austenite. . In addition, the Cr, Mo and V elements once dissolved by accelerated cooling also cause negative segregation to ferrite and positive segregation to austenite in this two-phase region retention, thereby reducing the yield ratio. The conflicting problem of ensuring strength can be solved. When the holding temperature in the two-phase region is lower than 730 ° C and higher than 830 ° C, the amount of reverse transformed austenite and ferrite is too small, and the yield ratio at the steel plate stage becomes high, and cold formed steel pipe The later yield ratio cannot satisfy the target yield ratio. The reason for quenching after maintaining the temperature in the two-phase region is to precipitate the bainite structure of the main phase and the island martensite phase by quenching from the reverse transformed austenite.

[Cold forming is performed in a range where t / d is 10% or less when the plate thickness is t (mm) and the outer cold bending diameter is d (mm)]
In the cold-formed steel pipe of the present invention, since the yield ratio on the tensile deformation side is 85% or less after processing, it has a cold-formed part where t / d is 10% or less. In order to form, cold forming is performed in a range where t / d is 10% or less.

  In the production method of the present invention, if necessary, (1) after reheating to a temperature range of 730 to 830 ° C. and then quenching, the steel sheet is tempered at 500 ° C. or less, and (2) the rolling is finished. Then, prior to accelerated cooling, it is preferable to add conditions such as online leveler correction, (3) cold forming at a steel plate temperature of 400 ° C. or lower. The reason for defining these requirements is as follows: It is as follows.

[After reheating to a temperature range of 730 to 830 ° C. and quenching, the steel sheet is tempered at 500 ° C. or lower]
It is also effective to selectively temper at 500 ° C. or less in order to eliminate the residual stress of the steel sheet subjected to the two-phase region quenching. When the tempering temperature at this time exceeds 500 ° C., C in the bainite structure generated as it is quenched is diffused and aggregated to generate pearlite, resulting in a decrease in strength. For these reasons, the tempering temperature is set to 500 ° C. or lower, preferably 480 ° C. or lower.

[Online leveler correction is performed prior to accelerated cooling after the rolling is completed]
Even after the end of rolling, even when a flat defect occurs at the front and rear ends of the rolled steel plate, the flatness is improved by hot straightening before direct quenching, so uniform cooling can be performed on the front and rear ends, and mechanical properties are improved. Stable and improved yield. For these reasons, it is effective to perform online leveler correction after the completion of rolling and before direct quenching.

[Cold forming at a steel plate temperature of 400 ° C. or lower]
As described above, the bending temperature (forming temperature) is acceptable not only at room temperature but also up to a temperature that does not impair the material of the steel sheet of the present invention (about 400 ° C.). In order to reduce the molding inhibition factor, it is also effective to perform selective molding (warm molding) at 400 ° C. or less at which the micro composition does not change and the dislocation density can be reduced. If the molding temperature at this time exceeds 400 ° C., C diffuses and a part of the bainite of the main phase starts to change to pearlite, resulting in a decrease in strength. A preferable temperature for this formation temperature is 300 ° C. or less.

  Hereinafter, the present invention will be described in more detail by way of examples.However, the present invention is not limited by the following examples as a matter of course, and may be implemented with modifications within a range that can meet the gist of the preceding and following descriptions. Of course, they are all possible and are included in the technical scope of the present invention.

  Steels having the chemical composition shown in Tables 1 and 2 were melted by a normal melting method, and any one of the following treatments was performed (types 1 to 3) to produce steel plates. Tables 1 and 2 also show the value of carbon equivalent Ceq, the value of [Mn] / [C] and the γ non-recrystallization temperature defined by the above formula (1).

[Processing procedure]
Type 1: After performing normal heating and hot rolling, direct quenching (DQ) is performed, and then quenching (Q ′) or 500 after two-phase region temperature (Ac ₁ point or more, less than Ac ₃ point) is maintained. Accelerated cooling was performed to below ℃.

Type 2: After completion of rolling, after slow cooling of about air cooled Ar less than ₃ points, were performed two-phase region temperature accelerated cooling or direct quenching from (Ar ₁ Tenkoe, Ar less than ₃ points) (DQ ').

  Type 3: After hot rolling, accelerated cooling and holding at a two-phase temperature were followed by accelerated cooling or direct quenching (DQ) again.

After that, some samples were tempered in addition to no tempering (T) at a temperature of less than ₁ Ac. The production conditions at this time are shown in Tables 3 to 5 below together with the value of the formula (2), the Ar ₃ transformation point, and the like.

About each obtained steel plate, t / d was changed and cold press molding was performed and the steel pipe was produced. While measuring the mechanical properties (yield point YP, tensile strength TS, uniform elongation δ _u ) of steel sheet and the type of microstructure, mechanical properties (yield point, tensile strength) in the tube axis direction (L direction) of the steel pipe TS, yield ratio YR and toughness), the amount of Cr, Mo, and V, and the microstructure were measured, and the materials were evaluated according to the following criteria.

[Solution amount of Cr, Mo, V]
About the solid solution amount of Cr, Mo, and V of a steel pipe, it calculated as the amount of each element deposited as the added amount of each element-precipitate. Regarding the element amounts of Cr, Mo, and V deposited as precipitates, the amount of precipitated elements was measured by an electrolytic extraction residue method in a cross section parallel to the surface of the outer t / 4 part of the steel pipe.

[Material Evaluation Criteria]
As the material evaluation criteria, the tensile strength TS in the tube axis direction of the steel pipe was set to 490 MPa or more, the yield ratio YR: 85% or less, and the fracture surface transition temperature (vTrs): −20 ° C. or less.

  The evaluation method of mechanical properties (steel plate and steel pipe), the toughness evaluation method of steel pipe, and the microstructure measurement method are as follows.

[Mechanical property evaluation method]
JIS Z 2201 No. 4 from t / 4 part (t is the plate thickness) of the steel plate to the L direction (rolling direction) and parallel to the tube axis of the outer t / 4 part of the steel pipe (corresponding to the main rolling direction of the steel plate) A specimen is taken and subjected to a tensile test according to JIS Z 2241. The mechanical properties of the steel sheet (yield point YP, tensile strength TS, uniform elongation δ _u ), and the mechanical properties of the steel pipe (yield point YP, tensile). The strength TS and the yield ratio (yield point / tensile strength × 100%: YR) were measured.

[Toughness evaluation method]
A specimen of JIS Z 2202 No. 4 was taken from the outer t / 4 part of the steel pipe in the direction parallel to the pipe axis (main rolling direction of the steel sheet), and Charpy impact test was conducted in accordance with JIS Z 2242. (VTrs) was measured.

[Microstructure measurement method]
At the steel sheet stage, the microstructure of the t / 4 part in the main rolling direction of the steel sheet is observed with an optical microscope, and the residual austenite γ _R present is subjected to X-ray diffraction of the steel sheet t / 4 part electropolished to 50 to 100 μm. The presence of residual austenite γ _R was confirmed from the peak intensity ratio between the α-Fe (200) plane and the γ-Fe (200) plane. Image analysis of the microstructure of the nital etched outer t / 4 part in the direction parallel to the pipe axis of the steel pipe (corresponding to the main rolling direction of the steel sheet) and t / 4 part in the main rolling direction of the steel sheet, Ferrite morphology, area fraction, bainite area fraction, and the like were measured. The island-like martensite phase was subjected to image analysis of a photograph of a microstructure obtained by etching 1/4 part of the plate thickness surface in the rolling direction with a repeller reagent, and the area fraction and the aspect ratio were measured.

  The steel pipes that satisfy the above-mentioned material standards were evaluated for weldability (weld crack resistance and HAZ toughness) by the following methods.

[Weld crack resistance]
In accordance with the y-type weld crack test method defined in JIS Z 3158, carbon dioxide gas welding was performed at a heat input of 1.7 KJ / mm, and the root crack prevention preheating temperature was measured. 25 degrees C or less was set as the pass.

[HAZ toughness]
Double-sided submerged arc welding (SAW) seam welding with 7KJ / mm heat input (X groove) was performed, and Charpy impact test piece (JIS Z 2204 No. 4) was taken from the outer t / 4 part in the direction perpendicular to the tube axis. The average impact absorption energy vE ₀ at 0 ° C. was determined (average value of three tests). The average vE ₀ was 47 J or more.

  The weldability test results are shown in the following Tables 6 to 8 together with mechanical properties (steel plate and steel pipe), microstructure and the like. From these results, it can be considered as follows. First, Experiment No. No. 1 is a control-rolled steel of V-added steel, and since Ceq exceeds the range specified in the present invention, the preheating temperature for preventing welding is as high as 50 ° C., and the HAZ toughness is also low.

  Experiment No. No. 2 is an accelerated cooling 450 ° C. stop material of Nb single added steel, and since polygonal ferrite is not generated in the microstructure, the yield ratio YR does not satisfy 85% or less of the target value after cold bending. ing.

  Experiment No. No. 3 is a steel obtained by quenching 450 ° C. of accelerated cooling of Nb-only added steel, followed by two-phase region temperature quenching (Q ′), and the area fraction of the polygonal ferrite phase is less than the range specified in the present invention. Therefore, the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  Experiment No. In the case of No. 14, although the value (A value) of the formula (2) is within the range specified by the present invention, the C content is larger than the range specified by the present invention. The toughness is low.

  Experiment No. No. 15, the Mn content is larger than the range defined in the present invention, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  Experiment No. In No. 18, the C content and the Mn content are less than the ranges defined in the present invention, and the tensile strength TS after cold bending does not satisfy the target value of 490 MPa or more.

  Experiment No. In No. 22, the Ti content is larger than the range defined in the present invention, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  Experiment No. In No. 24, the Mo content is larger than the range specified in the present invention, and the HAZ toughness does not satisfy the target value of 47 J or more.

  Experiment No. No. 29, the Si content is higher than the range specified in the present invention, the island martensite fraction of the steel pipe is higher than the range specified in the present invention, the yield ratio YR after cold bending However, it does not satisfy 85% or less of the target value.

  Experiment No. In No. 32, the Cu content is larger than the range defined in the present invention, and the preheating temperature for preventing welding is not satisfying the target of 25 ° C. or less. Further, the fracture surface transition temperature vTrs is −20 ° C. or lower, and the HAZ toughness does not satisfy the target value of 47 J or higher.

  Experiment No. In No. 35, the Nb content is larger than the range specified in the present invention, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  Experiment No. In No. 38, the content of Ca or REM is larger than the range specified in the present invention, and the fracture surface transition temperature vTrs after cold bending does not satisfy −20 ° C. or less.

  Experiment No. No. 51 has a heating temperature of 1300 ° C., and the fracture surface transition temperature vTrs after cold bending does not satisfy −20 ° C. or less.

  Experiment No. In No. 54, the heating temperature is 900 ° C., the cumulative rolling reduction is 100%, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value. In addition, Experiment No. 55, the cumulative rolling reduction is 80%, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  Experiment No. In No. 58, the accelerated cooling start temperature after rolling is 760 ° C., the polygonal ferrite fraction is 80 area%, and the tensile strength TS is reduced.

  Experiment No. In No. 60, the accelerated cooling stop temperature after rolling is 580 ° C., and the tensile strength TS is reduced.

  Experiment No. No. 61 has an accelerated cooling rate after rolling of 1.5 ° C./second, and the tensile strength TS is reduced.

  Experiment No. No. 63 has a reheating temperature before quenching of 850 ° C., a polygonal ferrite fraction of 35 area%, and the yield ratio YR after cold bending does not satisfy the target value of 85% or less. It has become.

  Experiment No. In No. 66, the reheating temperature before quenching is 700 ° C., and the tensile strength TS is lowered.

  Experiment No. In No. 68, the tempering temperature is 600 ° C., the polygonal ferrite fraction is 80 area%, the tensile strength TS is lowered, and the yield ratio YR after cold bending is 85% or less of the target value. Is not satisfied.

  Experiment No. No. 70 has a t / d of 15% during cold forming, and the yield ratio YR after cold bending does not satisfy 85% or less of the target value.

  On the other hand, all of the requirements specified in the present invention are satisfied (Experiment Nos. 4 to 13, 16, 17, 19 to 21, 23, 25 to 28, 30, 31, 34, 36, 37, 39). ˜50, 52, 53, 56, 57, 59, 62, 64, 65, 67, 69, 71) all satisfy the target values.

  Experiment No. The production points in 4 to 71 are as follows. That is, Experiment No. Nos. 4 to 38 are steels having the chemical composition shown in Tables 1 and 2 that were subjected to two-phase temperature quenching (Q ') after the end of rolling. No. 39 is further tempered (T).

  Experiment No. No. 40 was stopped at 450 ° C. after the end of rolling at the accelerated cooling. No. 41 is further tempered (T).

  Experiment No. No. 42 is a product which is slowly cooled (air-cooled) after the rolling, and directly quenched (DQ ') from the two-phase region temperature. No. 43 is further tempered (T).

Experiment No. 44, after the end of rolling, Ar ₁ Tenkoe, accelerated cooling to less than Ar ₃ point, and then held for 60 seconds in air, to produce a polygonal ferrite alpha _P, it is obtained by subsequently direct quenching.

Experiment No. 45, after the end of rolling, accelerated cooling to over Ar ₁ point and below Ar ₃ point, and then kept at the two-phase region temperature in an on-line holding furnace to generate polygonal ferrite α _P , and subsequently quenched. Yes, Experiment No. No. 46 is further tempered (T).

  Experiment No. 47 and 48 are obtained by changing t / d to 7.5, 5 (%) within the range defined by the present invention. Experiment No. 49 is a result of a plate thickness of 40 mm. Experiment No. 50 is press-bended after heating at 400 ° C.

  Experiment No. 51 to 54 are chemical components specified in the present invention and the heating temperature is changed in the range of 900 to 1300 ° C. Experiment No. 55 to 58 are chemical components defined in the present invention, and are obtained by changing the accelerated cooling start temperature after rolling.

  Experiment No. 59 and 60 are chemical components defined in the present invention, and the accelerated cooling stop temperature is changed. Experiment No. Reference numerals 60 to 62 are chemical components defined in the present invention, and change the accelerated cooling rate after rolling.

  Experiment No. 63 to 66 are chemical components defined in the present invention, and the heating temperature (Q ') during quenching is changed. Experiment No. 67 and 68 are chemical components specified in the present invention, and the tempering temperature (T) is changed.

  Experiment No. 69 is a chemical component specified in the present invention, and is obtained by performing on-line leveler correction before accelerated cooling. Experiment No. 70 is a chemical component defined in the present invention, and the cold bending t / d is outside the range defined in the present invention. Experiment No. 71 is a component specified in the present invention, and has a bending temperature of 400 ° C.

Claims (9)

C: 0.07 to 0.18% (meaning of mass%, hereinafter the same), Si: 0.05 to 1.0%, Mn: 0.7 to 1.7% (however, Mn content [Mn] And C content [C] ratio [Mn] / [C] ≦ 15), Ti: 0.002 to 0.025%, sol. In addition to containing Al: 0.005 to 0.1% and N: 0.001 to 0.008%, Cr: 0.6% or less (including 0%), Mo: 0.5% or less (0 %) And V: 0.08% or less (including 0%), including one or more selected from the group consisting of carbon equivalent Ceq values of 0.34 to It is within the range of 0.42%, and the A value shown by the following formula (2) satisfies 1.1 to 2.6, and the balance is made of a steel plate having a chemical composition composed of Fe and inevitable impurities. And the microstructure of the said steel plate is 40-70 area% polygonal ferrite phase, 0-20 area% pseudopolygonal ferrite phase, and 0-5 area%, and the aspect ratio (major axis / minor axis) is 4. 0.0 or less island martensite phase, balance is composed of bainite phase, 490 MPa class excellent in weldability, characterized in that it has a cold forming portion where t / d is 10% or less, where t is the t (mm) and the outer cold bending diameter is d (mm). Low yield ratio cold formed steel pipe.
Ceq = [C] + [Si] / 24 + [Mn] / 6 + [Ni] / 40 + [Cr] / 5
+ [Mo] / 4 + [V]) / 15 (1)
However, [C], [Si], [Mn], [Ni], [Cr], [Mo] and [V] are the contents (mass of C, Si, Mn, Ni, Cr, Mo and V, respectively). %).
A = (2.16 {Cr} +1) × (3.0 {Mo} +1) × (1.75 {V} +1) (2)
However, {Cr}, {Mo} and {V} indicate the amount of solid solution (mass%) in the steel sheet in Cr, Mo and V, respectively.   The 490 MPa class low yield ratio according to claim 1, further comprising Cu: 0.5% or less (not including 0%) and / or Ni: 3.0% or less (not including 0%). Cold formed steel pipe.   The 490 MPa class low yield ratio cold formed steel pipe according to claim 1 or 2, further comprising Nb: 0.015% or less (not including 0%).   The 490 MPa class low yield ratio cold formed steel pipe according to any one of claims 1 to 3, further comprising Ca: 0.005% or less (not including 0%).   The 490 MPa class low yield ratio cold-formed steel pipe according to any one of claims 1 to 4, further comprising rare earth element: 0.02% or less (not including 0%). In producing the cold-formed steel pipe according to any one of claims 1 to 5, the slab is heated to a temperature range of 950 to 1250 ° C, and the austenite non-recrystallization temperature A represented by the following formula (3): _After rolling to a steel plate by setting the cumulative reduction rate at _γ (° C.) or less to 60% or less (including 0%) to 4 ° C./second from the temperature above the Ar ₃ transformation point to 450 ° C. or less. Excellent weldability, characterized by accelerated cooling at a cooling rate, followed by reheating to a temperature range of 730 to 830 ° C. and quenching, followed by cold forming with a t / d of 10% or less. 490 MPa class low yield ratio cold-formed steel pipe manufacturing method.
A _γ (° C.) = 887 + 467 [C] + (6445 [Nb] −644√ [Nb]) +
(732 [V] -230√ [V]) + 890 [Ti] +363 [Al] -357 [Si]
... (3)
However, [C], [Nb], [V], [Ti], [Al], and [Si] indicate the contents (mass%) of C, Nb, V, Ti, Al, and Si, respectively.   The manufacturing method according to claim 6, wherein the steel sheet is tempered at 500 ° C. or lower after being reheated to a temperature range of 730 to 830 ° C. and then quenched.   The manufacturing method according to claim 6 or 7, wherein after the rolling is finished, online leveler correction is performed prior to accelerated cooling.   The manufacturing method according to any one of claims 6 to 8, wherein cold forming is performed at a steel plate temperature of 400 ° C or lower.