KR101069995B1 - High Strength Steel Sheet for Line-pipe and Manufacturing Method Thereof - Google Patents

Abstract

In the present invention, by weight%, C: 0.04 to 0.10%, Si: 0.05 to 0.50%, Mn: 1.4 to 2.0%, Al: 0.01 to 0.05%, P: 0.02% or less, S: 0.005% or less, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, Nb: 0.02 to 0.07%, V: 0.08% or less, Ni: 0.3% or less, Ca: 0.0005 to 0.004%, Cr: 0.1 to 0.3%, and Mo: Provides a steel sheet for high-strength line pipe containing at least one of 0.1 ~ 0.3%, the balance Fe and other inevitable impurities.

The steel sheet is mainly composed of a mixed structure of equiaxed ferrite and acicular ferrite, and has a high yield strength and tensile strength, including 1.5-7% of MA (Martensite / Austenite) structure, even when the coating heat treatment is performed after the pipe is manufactured to have excellent uniform elongation. It is possible to provide a high-strength line pipe steel sheet that can exhibit a continuous yield phenomenon due to the earthquake and the soil movement due to earthquake and frozen sea ice.

Steel plate for line pipe, equiaxed ferrite, needle ferrite, coating heat treatment, MA

Description

High Strength Steel Sheet for Line-pipe and Manufacturing Method Thereof}

The present invention relates to a method for manufacturing a high strength line pipe steel sheet used in pipeline construction, and more particularly, by adding Cr and Mo alone or in combination, to optimize the hot rolling process and the cooling after rolling to optimize corrosion resistance coating. The present invention relates to a method for producing a steel sheet for yield strength 80 ksi line pipe having improved plastic deformation and low temperature toughness even after heat treatment.

Recently, as the development of cold regions such as Siberia and Alaska, which are poor in climatic conditions, is becoming more active, the low-temperature toughness requirements for line pipe steel sheets are intensifying. In particular, in order to secure a sufficient stability of the pipe structure, in the past required a guarantee at 0 ℃ or -10 ℃ level, in recent years, a steel sheet of the level capable of guarantee -20 ℃ in the state made of pipes is required.

In general, in order for the line pipe steel sheet to be used safely at low temperatures, the DWTT characteristics exhibiting brittle fracture stopping characteristics should be excellent. In the conventional environment, it is possible to use the DWTT flexible wavefront at a pipe state of -10 ° C or more at 85% or more, but in an extreme cold environment, a steel sheet having a DWTT flexible wavefront ratio of -20 ° C or more at 85% is required.

In addition, in terms of the stability of the pipeline to earthquakes and the risk of collapse of the pipeline due to the movement of the soil during thawing of the freezing country, cold plastic deformation in the pipe axis direction is also required. It is known that the plastic deformation ability in the pipe axial direction is determined by the yield phenomenon and the uniform elongation in the pipe state, and in order to have resistance to earthquake and soil movement, the yield strength of 80 ksi grade is higher than 6% in the pipe state. Uniform elongation is required.

Conventionally, there have been studies on steel sheets having excellent plastic deformation ability in a pipe state, and representative examples thereof include Japanese Patent Laid-Open Publication Nos. 2003-293075, 2004-131799, and 2004-143499.

In Japanese Patent Laid-Open No. 2003-293075, an alloy element is controlled, and in particular, a steel that satisfies a relationship of Mn / 20 + Cu / 20 + Ni / 60 + Cr / 32 + Mo / 7 = 0.11, and Aspect ratio = The present invention relates to a steel member having excellent buckling-resistance characteristics using steel, which is composed of bainite structure containing 5 to 20% of MA (Martensite Austenite) coexistence structure of 4.0.

Further, Japanese Patent Laid-Open No. 2004-131799 is a bainite main structure in which the structure of the base material portion contains 20% or more of ferrite having an average particle diameter of 5 µm or less in terms of area ratio, and coarse austenite in the weld heat affected zone. The present invention relates to a high-strength steel pipe having excellent deformation performance, low temperature toughness and HAZ toughness, and a method for producing the same, which contain 5% or more of lower bainite in an area ratio in the tissue formed at the grain boundary.

Japanese Patent Application Laid-Open No. 2004-143499 has a microstructure containing ferrite with an average grain size of 10 μm or less and an area ratio of 70-90%, and contains 5-15% of retained austenite content when measured by X-ray diffraction. It relates to a high-strength steel pipe having excellent buckling characteristics, the slab is reheated at 1050 ℃ or more, and then rolled at 65% or more cumulative reduction ratio at an Ar3 temperature or more and 900 ° C or less, and 5 ° C from a temperature above the Ar3 temperature. It proposes a method of manufacturing steel pipe by cooling above / s and controlling the cooling conditions.

The inventions described in the above-mentioned Japanese Patent Laid-Open Publications are all related to pipes having excellent plastic deformation characteristics, and exhibit a uniform elongation of 6% or more. However, when the pipes are heat-treated at a maximum of 250 ° C. for corrosion-resistant coating on the pipes, Since discontinuous yielding occurs, plastic deformation capacity is inevitably deteriorated, so it is difficult to be easily applied to extreme regions or earthquakes.

Therefore, in order to solve this problem, a study has been conducted on a line pipe steel sheet exhibiting excellent plastic deformation ability even after performing a corrosion-resistant coating heat treatment. As such a conventional technique, Japanese Patent Laid-Open Publication No. 2002-220634, Korean Patent Patent Publication No. 2008-18285, and the like.

Japanese Laid-Open Patent Publication No. 2002-220634 proposes a method of limiting the contents of C and N, which are the causes of strain aging, and also suppressing the strain aging by adding Nb and Ti to combine these elements with C and N. do. However, since extremely low temperature rolling is performed for the precipitation of carbonitride during rolling, there is a problem in that productivity is extremely reduced and manufacturing costs are increased.

On the contrary, Korean Laid-Open Patent Publication No. 2008-18285 contains C: 0.03 to 0.1%, Si: 0.01 to 0.5%, Mn: 1.2 to 2.5%, and Al: 0.08% or less, and the metal structure contains ferrite and bay. It is a three-phase structure of knight and MA (Martensite / Austenite), and the area fraction of MA is 3 ~ 20%, and the online heat treatment using accelerated cooling and HOP (Heat treatment Online Process) is carried out on the steel sheet in which complex carbide is deposited on ferrite. It provides a method of manufacturing through. Since the prior art has to perform on-line heat treatment, there is still a problem that the process is complicated and productivity may be greatly reduced.

Therefore, it is easy to apply to the actual process in consideration of the additional equipment, productivity, etc. without adding a large amount of alloying elements to solve the above problems, it is necessary to provide a line pipe steel sheet having excellent plastic deformation ability even after corrosion-resistant coating have.

In the present invention, by weight%, C: 0.04 to 0.10%, Si: 0.05 to 0.50%, Mn: 1.4 to 2.0%, Al: 0.01 to 0.05%, P: 0.02% or less, S: 0.005% or less, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, Nb: 0.02 to 0.07%, V: 0.08% or less, Ni: 0.3% or less, Ca: 0.0005 to 0.004%, Cr: 0.1 to 0.3%, and Mo: Provides a steel sheet for high-strength line pipe containing at least one of 0.1 ~ 0.3%, the balance Fe and other inevitable impurities. The steel sheet has a mixed structure of equiaxed ferrite and acicular ferrite as a main phase, and contains 1.5-7% of MA (Martensite / Austenite) structure.

Furthermore, in the present invention, by weight%, C: 0.04 to 0.10%, Si: 0.05 to 0.50%, Mn: 1.4 to 2.0%, Al: 0.01 to 0.05%, P: 0.02% or less, S: 0.005% or less, Ti: 0.005-0.02%, N: 0.002-0.01%, Nb: 0.02-0.07%, V: 0.08% or less, Ni: 0.3% or less, Ca: 0.0005-0.004%, Cr: 0.1-0.3% and Mo: slab containing at least one of 0.1 to 0.3%, and the remaining Fe and other unavoidable impurities, heating at 1100 ~ 1200 ℃, at an average rolling reduction of 10% or more per pass at a temperature of 1050 ℃ or more Provided is a method for producing a high strength linepipe steel sheet comprising recrystallization rolling, hot rolling with a cumulative reduction of more than 70% in the temperature range of 950 ° C to Ar3, and quenching to a very low cooling stop temperature, that is, Ms-60 ° C followed by air cooling. do.

The steel sheet according to the present invention has a high yield strength and tensile strength, but can exhibit a continuous yield phenomenon due to excellent uniform elongation even after coating heat treatment after pipe manufacturing, high strength line that can withstand soil movement due to earthquake and thawing of soil It is possible to provide a steel sheet for pipes.

The inventors of the present invention have conducted a study on the production of line pipe steel sheet having excellent yield strength. As a result, the steel sheet is rolled in the austenite recrystallization and uncrystallization region, and Ms temperature (° C), which is the theoretical martensite transformation temperature, is 561 to 474 *. When cold-cooled accelerated cooling is performed below (% C)-33 * (% Mn)-17 * (% Ni)-17 * (% Cr)-21 * (% Mo), the austenite effective grain size is remarkably refined and DWTT The properties were greatly improved, and it was confirmed that the equiaxed ferrite, acicular ferrite, and M / A structure were finely and uniformly distributed so that the plastic deformation ability was excellent even after the corrosion coating, and thus the present invention was completed based on the experimental results.

Hereinafter, the present invention will be described in detail. However, the content% of the component system below means weight%.

C: 0.04-0.10%

C is the most economical and effective element for reinforcing steel, but when added in a large amount in excess of 0.10%, C is not good in weldability, formability, toughness, etc., and therefore, the upper limit thereof is limited to 0.10% in consideration of this. However, if the amount of C added is less than 0.04%, other expensive alloying elements such as Mo, Ni, etc. must be added in large amounts in order to exhibit the same strength, which is not preferable because it is very uneconomical. Therefore, the content of C in the present invention is limited to 0.04 ~ 0.10%.

Si: 0.05-0.50%

Since Si acts to deoxidize molten steel, it is useful for obtaining a deoxidation effect. Since Si can be used as a solid solution strengthening element, Si is added in an amount of 0.05 to 0.50%. If the amount of Si added is less than 0.05%, the deoxidation of molten steel may not be sufficient, so the toughness may be reduced. On the other hand, if the amount of Si is added more than 0.5%, the red scale by Si may be formed during hot rolling, which may lower the surface quality of the steel sheet. And adversely affect the weld toughness.

Mn: 1.4-2.0%

Mn is an effective element to solidify steel, and its content must be more than 1.4% to show sufficient yield strength (high strength required in 80ksi steels) with the effect of increasing hardenability. However, if the Mn content is added in excess of 2.0%, it is not preferable because the segregation part is greatly developed at the center of thickness when casting the slab in the steelmaking process, and the weldability of the final product is deteriorated.

Al: 0.01 ~ 0.05%

Al is added as a deoxidizer along with Si in steelmaking and is added at least 0.01% because it has a solid solution strengthening effect. However, when Al is added in excess of 0.05%, impact toughness is lowered, and if less than 0.01%, deoxidation effect may be insufficient and the toughness may be lowered. In the present invention, Al is added in the range of 0.01 to 0.05%.

P: 0.02% or less

P is an impurity element existing in steel, and it is preferable to reduce it as much as possible because it segregates in the center of the steel sheet and impairs toughness. However, since the time and cost are required to manage too low, the upper limit is 0.02%.

S: 0.005% or less

S is also an impurity element present in steel, and is preferably combined with Mn or the like to form a non-metallic inclusion, thereby greatly reducing the toughness and strength of the steel, and thus reducing it as much as possible. In particular, in order to secure brittle fracture stop characteristics at cryogenic temperatures, the upper limit is set at 0.005%.

Nb: 0.02-0.07%

Nb is added to 0.02% or more in the present invention because Nb is a very useful element for refining grains and at the same time serves to promote formation of acicular ferrite or bainite, which is a high-strength structure, to greatly improve the strength of steel. However, if Nb is added in excess of 0.07%, there is a problem that the weldability may be lowered. Therefore, the content of Nb is limited to 0.02 ~ 0.07%.

V: 0.08% or less

V is an element useful for improving the strength by enhancing the precipitation by V (C, N) precipitates and improving the quenching property. However, when a large amount is added, weldability and toughness deteriorate, so the amount of V is limited to 0.08% or less.

Ti: 0.005-0.02%

Ti is an element that forms a TiN precipitate during the solidification process of steel, suppresses austenite grains during slab heating and hot rolling, thereby minimizing the final grain size, thereby improving the toughness of the steel. If the Ti content is less than 0.005%, the formation of TiN precipitates is insufficient, and the effect of inhibiting particle size growth cannot be expected.In contrast, if the Ti content exceeds 0.02%, the TiN precipitates coarsely when the slab is heated due to the excessive presence of solute Ti. Not suitable for miniaturization Therefore, in the present invention, the Ti content is limited to 0.005 to 0.02%.

N: 0.002-0.01%

The reason for component limitation of N is attributable to the above Ti addition. In general, N is dissolved in the steel and precipitates to increase the strength of the steel, which is much greater than carbon. However, N present in steel is known to reduce toughness, and it is a general trend to reduce nitrogen content as much as possible. However, in the present invention, it is not desirable to excessively reduce N to less than 0.002%, since the proper amount of N is present to react with Ti to form TiN and to inhibit grain growth during reheating. In addition, when N is too small, the content of TiN precipitates is small, so that the particle size and growth inhibitory effect are not good. On the contrary, when there is too much N content, since N exists in solid-solution N instead of TiN form, toughness may fall significantly, the upper limit shall be 0.01%.

Ni: 0.3% or less

Ni is an element used to improve toughness, and in the present invention, it also contributes to the action of improving the brittle fracture stopping property and the improvement of strength. Ni increases toughness as the amount added increases, but it is expensive and toughness does not increase proportionally as the amount added increases, so it is limited to 0.3% or less.

Mo: 0.1 ~ 0.3%

Mo is very effective in increasing the strength of the material, and promotes the formation of bainite, which is a low temperature transformation structure, and helps to obtain high strength and high toughness at the same time. In addition, since the element contributes to the formation of the MA structure during the cold-accelerated cooling of the element having a large influence on the hardenability improvement, it improves the plastic deformation capacity of the pipe, which affects the resistance to pipe deformation in the extreme regions. Therefore, in the present invention, 0.1% or more of Mo is added. However, Mo is an expensive element, and if the amount is increased, weldability is lowered, so the upper limit thereof is limited to 0.3%.

Cr: 0.1-0.3%

Cr is effective to increase the strength of the material like Mo. In addition, there is a function to promote the formation of MA during cold-cooled accelerated cooling, it is possible to replace the expensive Mo element. In order to obtain such an effect, in the present invention, 0.1% or more of Cr is added. However, excessive addition of Cr may cause deterioration of weldability, so the upper limit is 0.3%.

Ca: 0.0005-0.004%

Ca is an element that spheroidizes the MnS inclusions and suppresses crack formation around the inclusions, and is added at least 0.0005%. If it is less than 0.0005%, the inclusion spheroidizing effect does not appear, whereas if it exceeds 0.004%, impact toughness may be lowered by the formation of a large amount of CaO-based inclusions.

The slab including the above-described component system is made of a steel sheet by the manufacturing method of the present invention. Hereinafter, the steel sheet manufacturing method of the present invention will be described in more detail.

In the present invention, the slab is first heated to 1100 ~ 1200 ℃. And recrystallization rolling is made at 10% or more per pass at 1050 ° C or higher and cooled to 950 ° C or lower. Thereafter, hot rolling is performed at a cumulative reduction of 70% or more in the temperature range of 950 ° C to Ar3, followed by water cooling at a cooling rate of 10 ° C / s or more, and thus Ms temperature [Martensite transformation start temperature, 561-474 * (% C)- 33 * (% Mn) -17 * (% Ni) -17 * (% Cr) -21 * (% Mo)] after finishing water cooling at a temperature of 60 ° C. or lower, and then cooling the plate to room temperature Is manufactured.

Slab heating temperature: 1100 ~ 1200 ℃

The temperature of reheating the slab has an important meaning in the present invention. The high strength and high toughness properties in the present invention result from the presence of fine acicular ferrite formed by the addition effect of fine carbonitride precipitates, Cr, Mo and solid solution Nb. Therefore, it is necessary to heat the slab to 1100 ° C. or higher to dissolve NbC before hot rolling so that Nb is present in an atomic state. However, since the coarse TiN precipitation may occur upon reheating when the heating temperature exceeds 1200 ° C, the slab heating temperature in the present invention is 1100 to 1200 ° C.

Recrystallization rolling conditions: average rolling reduction of 10% or more per pass above 1050 ° C

The heated slabs are recrystallized at 1050 ° C. or higher. In the present invention, since the content of Nb is 0.02 ~ 0.07%, austenite is completely recrystallized when rolling at a temperature of 1050 ℃ or more.

Therefore, rolling at a temperature below 1050 ° C. (exactly 950-1050 ° C.), which is a partial recrystallization zone of austenite, results in a coarse structure of coarse austenite and fine austenite, resulting in a partially coarse grain in the final structure. Toughness can be greatly reduced.

Therefore, it is necessary to completely recrystallize austenite grains by rolling at an average rolling reduction of 10% or more per pass above 1050 ° C. If the average rolling reduction per pass is less than 10%, the recrystallized austenite grain size may be coarse and the toughness may deteriorate, so the average rolling reduction of the recrystallized rolling is made 10% or more.

Cooling condition after recrystallization rolling: Cooling to 950 ℃ or lower

In the present invention, after recrystallization rolling, the slab is cooled by air without rolling to 950 ° C or lower. This is because rolling in this section can cause partial recrystallization, which increases the likelihood of brittle fracture due to coarse austenite grain size.

Unrecrystallized rolling conditions: Temperature range from 950 ° C to Ar3, cumulative reduction of more than 70%

The slab cooled to 950 ° C. or lower is hot rolled with a cumulative reduction of 70% or more in the temperature range of 950 ° C. to Ar 3, which is a non-recrystallized region. The non-recrystallized rolling process is for obtaining fine needle-like ferrites and equiaxed ferrites by inducing grain boundaries and intragranular deformed structures of elongated austenite, and greatly improving the strength, plastic deformation ability, and brittle fracture stopping properties.

The larger the cumulative rolling reduction, the more effective for improving toughness. If the cumulative rolling reduction is less than 70%, the brittle fracture resistance is not sufficiently obtained, so the cumulative rolling reduction in the present invention is 70% or more.

The lower the hot rolling finish temperature, the higher the austenite strain is, which is effective in improving toughness, but below the Ar3 temperature, the toughness may be deteriorated by the formation of the deformed equiaxed ferrite structure during rolling and is not good for plastic deformation. The minimum temperature of rolling is made into Ar3 temperature. If there is a strained equiaxed ferrite formed by rolling in a section below the Ar3 temperature, strain aging, a phenomenon in which dislocations are adhered to solute carbon, may occur, and thus plastic deformation of the steel may be drastically reduced.

Cooling conditions: water cooled to 10 ° C / s or more, up to 60 ° C or less relative to Ms temperature

After the hot rolling is finished, the steel sheet is cooled by quenching, preferably by water cooling, to a temperature of 60 ° C. or lower relative to the Ms temperature (Martensite transformation start temperature). At this time, when the cooling rate is low, the amount of the equiaxed ferrite with low strength may be excessively large and the grains may grow to have a coarse particle size, which is not good for strength or toughness. Therefore, it is preferable to make cooling rate into 10 degreeC / s or more. In addition, if the cooling rate is high, the low-temperature transformation tissue may be generated, but when controlling the component system as in the present invention, even if the cooling rate to some extent it can be seen that the low-temperature transformation tissue such as martensite is not formed, Figure 3 and As shown in FIG. 4, only acicular ferrite and / or polygonal ferrite are formed, thereby securing toughness.

In addition, the cooling stop temperature has a particularly important meaning not only in terms of strength-toughness of steel, but also in terms of securing plastic deformation after corrosion-resistant coating heat treatment. As a result of the experiments of the present inventors, in the present invention, the cooling stop temperature was limited very low, because it was found that the plastic deformation capacity after the corrosion-resistant coating changed significantly based on the Ms-60 ° C. If the cooling stop temperature is higher than Ms-60 ° C, the MA content present in the steel is less, resulting in discontinuous yielding after corrosion-resistant coating heat treatment and lowering of plastic deformation.

On the other hand, if the cooling stop is made at Ms-60 ℃ or less as shown in the conditions of the present invention, since the MA fraction is 1.5% or more in the steel, the strain aging after corrosion-resistant coating is suppressed, the continuous yield occurs, the uniform elongation is excellent, and the plastic strain is Significantly improved. This effect is considered to be because, when MA is 1.5% or more, the movable potential increases locally around the MA phase, and C cannot fix all of the movable potentials. Therefore, as the MA content increases, the movable potential in the steel increases and the strain aging suppression after the corrosion resistant coating heat treatment becomes easy. However, excessively increasing the MA fraction may cause a problem of deterioration of toughness, so the upper limit of the MA content is limited to 7%.

Hereinafter, the present invention will be described in more detail with reference to the following examples.

(Example)

After preparing slabs including alloy elements as shown in Table 1, each slab was hot rolled to prepare a steel plate having a thickness of 25 mm.

Steel grade C Si Mn P S Nb V Ni Cr Mo Al Ca Ti N Invention
(A) 0.05 0.25 1.80 0.010 0.001 0.040 - 0.25 - 0.25 0.03 0.0010 0.010 0.004 Invention
(B) 0.06 0.18 1.80 0.007 0.002 0.040 - 0.25 0.1 0.15 0.03 0.0006 0.010 0.0043 Invention
(C) 0.07 0.20 1.80 0.010 0.002 0.40 - 0.15 0.1 0.15 0.03 0.0010 0.015 0.0040 Invention
(D) 0.075 0.31 1.57 0.004 0.001 0.048 0.06 0.15 0.2 - 0.05 0.0005 0.007 0.0049 Comparative material
(E) 0.070 0.24 1.55 0.010 0.002 0.042 0.05 0.19 - 0.08 0.03 0.0011 0.015 0.0030 Comparative material
(F) 0.070 0.22 1.55 0.010 0.001 0.040 - 0.15 0.08 - 0.03 0.0010 0.015 0.0034 Comparative material
(G) 0.051 0.27 1.24 0.008 0.002 0.037 0.05 0.19 - 0.20 0.03 0.0007 0.015 0.0050

In addition, each invention material (A ~ D) and the comparative material (E ~ G) was prepared in a steel sheet under the production conditions shown in Table 2, and subjected to tensile, impact and DWTT test. In addition, a steel pipe having a diameter of 48 inches is manufactured for the produced steel sheet, and a tensile test piece of ASTM sub size standard is taken in the longitudinal direction of the steel pipe, followed by a tensile test after heat treatment at a maximum temperature of coating heat treatment at 250 ° C. Behavior and uniform elongation were measured. The experimental results are shown in Table 3 below.

Psalter Steel grade Slab heating temperature
(℃) Crude rolling
End
Temperature
(℃) Finish rolling
start
Temperature
(℃) Rolling
End
Temperature
(℃) Unresolved station
accumulate
Rolling reduction
(%) Cooling
start
Temperature
(℃) Cooling
End
Temperature
(℃) theory
Ms temperature-60 ℃ * Cooling
speed
(℃ / s) Inventive Example One A 1150 1070 900 780 74 720 348 408 16 2 B 1130 1050 930 820 79 740 301 404 35 3 B 1130 1050 950 840 80 760 381 404 34 4 C 1130 1060 940 840 75 740 280 401 28 5 C 1120 1050 930 830 75 730 380 401 23 6 D 1130 1060 930 830 75 730 343 407 32 7 D 1130 1060 930 825 75 730 394 407 28 Comparative example One A 1150 1060 920 770 71 700 550 408 26 2 A 1150 1060 920 760 75 720 480 408 20 3 B 1130 1050 930 810 79 730 450 404 24 4 B 1130 1050 950 840 80 760 418 404 30 5 C 1120 1050 920 760 75 725 467 401 21 6 D 1140 1060 920 820 75 730 414 407 26 7 D 1140 1060 920 760 75 720 390 407 8 8 D 1140 1060 930 770 60 720 390 407 28 9 D 1150 1060 990 880 75 800 380 407 30 10 E 1150 1060 950 800 79 740 406 410 28 11 E 1150 1060 940 790 75 730 351 410 32 12 F 1150 1060 950 840 75 760 343 411 35 13 F 1150 1060 920 820 75 760 314 411 36 14 G 1150 1060 950 790 75 720 310 430 28

* Ms temperature (° C) = 561-474 * (% C)-33 * (% Mn)-17 * (% Ni)-17 * (% Cr)-21 * (% Mo)

Psalm No Steel grade M / A
Fraction
(%) Rolling right angle
Direction yield strength
(MPa) Rolling right angle
Directional tensile
burglar
(MPa) Coating Elongation Uniform Elongation
(Rolling direction,%) Coated Copying Materials
Yield behavior
(Rolling direction) -20 ℃ DWTT SA% Inventive Example One A 2.1 580 700 7.9 Continuous surrender 99 2 B 2.6 592 765 7.4 Continuous surrender 95 3 B 1.6 609 708 8.4 Continuous surrender 99 4 C 3.6 557 662 7.8 Continuous surrender 99 5 C 1.7 567 651 7.5 Continuous surrender 99 6 D 2.3 615 727 6.2 Continuous surrender 96 7 D 1.8 603 714 6.5 Continuous surrender 97 Comparative example One A 0.3 559 689 8.9 Discontinuous surrender 99 2 A 0.5 592 695 7.6 Discontinuous surrender 90 3 B 0.8 604 680 7.7 Discontinuous surrender 97 4 B 0.7 602 690 8.1 Discontinuous surrender 97 5 C 1.2 571 636 8.5 Discontinuous surrender 99 6 D 1.1 586 683 6.7 Discontinuous surrender 96 7 D 1.5 449 692 5.9 Continuous surrender 98 8 D 2.2 576 710 5.8 Continuous surrender 70 9 D 2.3 574 695 5.8 Continuous surrender 60 10 E 0.9 589 693 6.4 Discontinuous surrender 80 11 E 1.1 561 700 6.7 Discontinuous surrender 95 12 F 0.8 558 668 7.0 Discontinuous surrender 99 13 F 0.7 569 655 7.5 Discontinuous surrender 90 14 G 1.9 545 615 8.8 Continuous surrender 50

As shown in Table 3, Inventive Examples 1 to 7, which manufactured the inventive materials A to D satisfying the component range of the present invention under the manufacturing conditions of the present invention, yield strength 557 to 615 Mpa, tensile strength 651 to 765 Mpa, and coating heat treatment. After the continuous elongation at 6.2 ~ 8.4%, and the DWTT ductile fracture rate at -20 ℃ is 95 ~ 99%, the strength and toughness required for API-X80 grade low temperature line pipe steel In addition, it can be seen that it has a plastic deformation capacity that can withstand soil movement due to earthquake and freezing of ice even after corrosion-resistant coating heat treatment.

On the other hand, the alloy element composition of the present invention satisfies, but the accelerated cooling stop temperature is higher than the theoretical Ms temperature-60 ℃, API-X80 in the case of Comparative Examples 1 to 6 where the M / A content in the steel is less than 1.5% Yield strength and DWTT characteristics required for the grade were satisfied to some extent, but discontinuous yield occurred after the corrosion-resistant coating heat treatment, and the plastic deformation was greatly reduced.

1 and 2 show these experimental results. FIG. 1 is a stress-strain curve of tensile strength test for Inventive Example 2 and FIG. 2 for Comparative Example 3 after heat treatment at 250 ° C. for 5 minutes. 1 and 2 show that there is a remarkable difference in the tensile properties after heat-resistant coating heat treatment of the inventive and comparative pipes, respectively.

In addition, in FIG. 3 and FIG. 4, in order to clarify the reason why the difference in plastic deformation and low-temperature toughness appear between the invention example with a cooling stop temperature of Ms-60 degrees C or less, and the comparative example with a comparatively high cooling stop temperature, Inventive Example 2 and a comparison The photograph which observed Example 3 with the optical microscope is shown. Each of them was photographed by etching with picral solution for the appearance of MA (Martensite / Austenite) constituent. As can be seen in Figures 3 and 4, both the invention examples and the comparative examples are evenly distributed fine equiaxed ferrite and acicular ferrite having an average particle size of 10㎛ or less. However, the invention example shows a great difference in the MA fraction and distribution from the comparative example. That is, in the example of FIG. 3, the MA fraction was 2.5%, which is high, and showed a very fine distribution. In the comparative example, the MA fraction was less than 0.5%. From these results, it can be seen that the reason why the invention example produced at the low temperature cooling end temperature is excellent in strength-toughness and plastic deformation capacity is due to formation of fine MA effective for structure refinement and plastic deformation capacity by step-down cooling.

In Comparative Example 7 in which the cooling rate was out of the manufacturing conditions of the present invention, the yield strength was low, which did not satisfy the level of the yield strength 80 ksi grade steel, and the uniform elongation was less than 6%. In addition, in the case of Comparative Example 8 in which the unrecrystallized region rolling reduction did not satisfy the manufacturing conditions of the present invention, the uniform elongation was low and the DWTT characteristics were greatly deteriorated. Even in the case of Comparative Example 9 in which the unrecrystallized region rolling start temperature was higher than the production conditions of the present invention, the uniform elongation and the DWTT characteristics simultaneously decreased.

On the other hand, even if the manufacturing conditions of the present invention is satisfied, Comparative Examples 10 to 13 manufactured with comparative materials E and F having Cr and Mo contents lower than the inventive steel, despite the accelerated cooling stop temperature is Ms-60 ℃ or less, The M / A content was less than 1.5%, resulting in discontinuous yield after corrosion-resistant coating, and greatly reduced plastic deformation. Comparative material 14 using comparative material G having a low Mn content out of the range of components of the invention steel showed a low yield strength even if it satisfies the manufacturing conditions of the present invention. It is well understood that it is not suitable for use in steel.

1 is a stress-strain graph of the tensile test for Inventive Example 2 of the Examples of the present invention.

2 is a stress-strain graph of the tensile test for Comparative Example 3 of the Examples of the present invention.

Figure 3 is a photograph of the microstructure of Example 2 of the present invention observed with an optical microscope.

Figure 4 is a photograph of observing the microstructure of Comparative Example 3 of the embodiment of the present invention with an optical microscope.

Claims (5)

By weight%, C: 0.04 to 0.10%, Si: 0.05 to 0.50%, Mn: 1.4 to 2.0%, Al: 0.01 to 0.05%, P: 0.02% or less, S: 0.005% or less, Ti: 0.005 to 0.02% , N: 0.002-0.01%, Nb: 0.02-0.07%, V: 0.08% or less (excluding 0%), Ni: 0.3% or less (excluding 0%), Ca: 0.0005-0.004%, At least one of Cr: 0.1-0.3% and Mo: 0.1-0.3%, Balance includes Fe and other unavoidable impurities, A steel sheet for high-strength line pipe, comprising a mixed structure of equiaxed ferrite and acicular ferrite as a main phase, and containing 1.5 to 7 area% of MA (Martensite / Austenite) structure. delete By weight%, C: 0.04 to 0.10%, Si: 0.05 to 0.50%, Mn: 1.4 to 2.0%, Al: 0.01 to 0.05%, P: 0.02% or less, S: 0.005% or less, Ti: 0.005 to 0.02% , N: 0.002-0.01%, Nb: 0.02-0.07%, V: 0.08% or less (excluding 0%), Ni: 0.3% or less (excluding 0%), Ca: 0.0005-0.004%, Cr : At least one of 0.1 to 0.3% and Mo: 0.1 to 0.3%, the remainder is for slabs containing Fe and other unavoidable impurities, Heating step of heating at 1100 ~ 1200 ℃; A recrystallization rolling step of recrystallization rolling at an average rolling reduction of 10% or more per pass at a temperature of 1050 ° C or higher; A hot rolling step of hot rolling at a cumulative reduction of 70% or more in a temperature range of 950 ° C to Ar3; And Cooling step of quenching to below Ms-60 ° C. followed by air cooling Method for producing a high-strength line pipe steel sheet comprising a. The method of claim 3, wherein the recrystallization rolling step comprises cooling the steel sheet without rolling to 950 ° C. or less after rolling. 5. The method of claim 3, wherein the cooling rate during the quenching process is 10 ° C / s or more.

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