JP5151034B2 - Manufacturing method of steel plate for high tension line pipe and steel plate for high tension line pipe - Google Patents

Description

  The present invention relates to a method for manufacturing a steel sheet for high-strength line pipe, and more particularly to a method for manufacturing a steel sheet for high-strength line pipe excellent in strength, toughness, DWTT characteristics, and CTOD characteristics suitable for use in pipelines.

  In recent years, line pipes used for transportation of natural gas and crude oil have been strengthened year by year in order to improve transport efficiency by increasing pressure and to improve local welding efficiency by reducing wall thickness. So far, X100 grade line pipes have been put into practical use according to the API standard, and further, the demand for X120 grade exceeding tensile strength 900 MPa is being realized.

  For example, Patent Document 1 relates to a welded steel pipe for high-strength line pipes and a method for producing a high-strength thick steel plate as a raw material thereof, in order to obtain high strength and high toughness while reducing the amount of expensive alloy element addition. Discloses technology relating to accelerated cooling and tempering conditions.

Also, Patent Document 2 discloses a component design technique for the base metal that reduces the amount of alloying elements added as in Patent Document 1 and that provides high strength and high toughness in the weld metal of the longitudinal seam weld. Yes.
JP 2002-173710 A JP 2000-355729 A

  However, when increasing the strength by means such as accelerated cooling while keeping the amount of alloying elements in the base metal low, the cooling rate of the HAZ is constant unless the vertical seam welding method is appropriately selected. While the strength remains low, there is a difference between the base material and the strength of the weld heat affected zone (Heat Affected Zone, hereinafter abbreviated as HAZ).

  As a result, even if the strength of only the base metal part is increased, for example, when a real pipe test such as a water pressure test is performed, the HAZ part having a low strength is broken, and there remains a problem in safety in practical use. Increasing the strength of the weld metal in the longitudinal seam welded part serves to compensate for the lack of joint strength caused by such a HAZ softened part, but it cannot be said to be sufficient.

  On the other hand, it is widely known that the lower bainite structure is used as a microstructure structure excellent in the strength and toughness balance of the base material, and it is possible to obtain the lower bainite structure by appropriate component design including alloy addition. . By effectively utilizing this structure, the joint strength after welding can be ensured. However, the optimum manufacturing method for obtaining this structure has not been clarified so far.

  An object of the present invention is to provide a method for producing a steel sheet for a line pipe, which is a component system capable of utilizing the lower bainite structure and is excellent in strength / toughness balance of a base material, DWTT characteristics, and further CTOD characteristics.

  In order to achieve the above-mentioned problems, the present inventors diligently studied various factors affecting low-temperature toughness at a tensile strength of 900 MPa or more for steel sheets for line pipes.

  As a result, the following knowledge was obtained when quenching was stopped at a temperature below the Ms point, maintained for a short time in the temperature range, then air-cooled or maintained for a short time in the temperature range, and then rapidly heated and tempered. It was.

  1 Quenching at a specific cooling rate from the austenite temperature, stop quenching at a temperature range below the Ms point, maintain for a specific time in the temperature range, and then air-cooled. A mixed structure of sites and lower bainite is obtained, and the strength-toughness balance is improved.

In low carbon steel, austenite is transformed during quenching to form a lath structure. At this time, C is concentrated between the laths, and untransformed austenite (γ) film tends to remain, which is harmful to toughness. Although it is easy to form island-like martensite (hereinafter referred to as MA), quenching was stopped at a temperature below the Ms point, and after maintaining for a short time in the temperature range, it was generated during quenching. In addition to the tempering of martensite, the lower bainite is fine and abundantly produced, and MA is not reduced, thereby improving toughness.

  2 Quenching is stopped at a temperature below the Ms point, maintained for a short time in the temperature range, and then rapidly heated at a specific temperature increase rate to perform tempering, thereby reducing strength deterioration and improving toughness.

  Using equipment with a high frequency induction heating device on 3 lines, it is possible to stably give the above heat history, particularly quenching at a temperature below the Ms point and maintaining for a short time in the temperature range. Possible and extremely effective.

Furthermore, in addition to the above findings, it has been found that CTOD characteristics are improved by increasing the cumulative rolling reduction in the austenite non-recrystallization temperature region in hot rolling before quenching.
In the present invention, “high tension” and “high strength” are strengths having a tensile strength of 900 MPa or more.

The present invention has been completed based on such findings and further studies. That is, the gist of the present invention is as follows.
1. % By mass, C: 0.04 to 0.12%, Si: ≦ 0.50%, Mn: 1.80 to 2.50%, P ≦ 0.010%, S ≦ 0.002%, Al: 0.01-0.08%, Cu: 0.01-0.8%, Ni: 0.1-1.0%, Cr: 0.01-0.8%, Mo: 0.01-0. 8%, Nb: 0.01-0.08%, V: 0.01-0.10%, Ti: 0.005-0.025%, B: 0.0005-0.0030%, Ca: ≦ 0.01%, REM: ≦ 0.02%, N: 0.001 to 0.006%, and the steel consisting of the balance Fe and inevitable impurities is heated to 1000 to 1200 ° C. in Ar ₃ transformation point or higher, the cumulative rolling reduction in the austenite non-recrystallization temperature region subjected to hot rolling 50% or more, then, the temperature range of not lower than Ar ₃ transformation point, Quenching cooling in a temperature range of 300 ° C. or more below the martensite transformation start temperature (Ms point) defined by the following formula (2) at a cooling rate of the martensite formation critical cooling rate Vcrm defined by the formula (1) The tensile strength is characterized in that after cooling to the stop temperature, the steel temperature is maintained within the cooling stop temperature ± 50 ° C. by online heating for 60 s to 300 s after the cooling stop, and then air cooling to room temperature. A method for producing a steel sheet for a high-strength line pipe having a characteristic of 900 MPa or more, exceeding 200 J in a Charpy impact test at −30 ° C., and having a ductile fracture surface ratio in a DWTT test of 85% or more at −20 ° C.
logVcrm = 2.94-0.75β
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo)
... (1)
Here, Vcrm: Martensite formation critical cooling rate (° C./s)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
... (2)
2. % By mass, C: 0.04 to 0.12%, Si: ≦ 0.50%, Mn: 1.80 to 2.50%, P ≦ 0.010%, S ≦ 0.002%, Al: 0.01-0.08%, Cu: 0.
01-0.8%, Ni: 0.1-1.0%, Cr: 0.01-0.8%, Mo: 0.01
-0.8%, Nb: 0.01-0.08%, V: 0.01-0.10%, Ti: 0.00
5 to 0.025%, B: 0.0005 to 0.0030%, Ca: ≦ 0.01%, REM:
≦ 0.02%, N: 0.001 to 0.006%, and the steel consisting of the remainder Fe and inevitable impurities is heated to 1000 to 1200 ° C., and the rolling end temperature is at least the Ar ₃ transformation point . Hot rolling with a cumulative reduction ratio of 50% or more in the austenite non-recrystallization temperature range, and then a martensite formation critical cooling rate Vcrm defined by the following formula (1) from a temperature range of Ar ₃ transformation point or higher After cooling to the quenching cooling stop temperature in the temperature range of 300 ° C. or higher, which is equal to or lower than the martensite transformation start temperature (Ms point) defined by the following formula (2), for 60 s to 300 s after the cooling stop, the online heating, the temperature of the steel is maintained within the cooling stop temperature ± 50 ° C., then, to rapid heating immediately from the temperature to 450 ° C. or higher Ac ₁ transformation point of the temperature range at 1 ° C. / s or more heating rate Features and tensile strength than 900MPa that, in more than 200J In Charpy impact test at -30 ° C., the ductility fracture rate at -20 ° C. In DWTT test for high tension line pipe having a 85% or more properties A method of manufacturing a steel sheet.
logVcrm = 2.94-0.75β
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo)
... (1)
Here, Vcrm: Martensite formation critical cooling rate (° C./s)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
... (2)
3 . Heating after water cooling is performed by a high frequency induction heating device installed on the same line as the cooling device on the downstream side of the cooling device, wherein the tensile strength according to 1 or 2 is 900 MPa or more and −30 A method for producing a steel sheet for high-strength line pipes having a characteristic that the Charpy impact test at 200 ° C exceeds 200 J and the ductile fracture surface ratio at -20 ° C in the DWTT test is 85% or more .
4 . It is manufactured by the manufacturing method according to any one of 1 to 3 , and the microstructure is a mixed structure of tempered martensite and lower bainite having a volume ratio of 90% or more, and further contains at least 50% or more of the lower bainite. A steel plate for high-strength line pipes having a tensile strength of 900 MPa or more, a property of exceeding 200 J in a Charpy impact test at −30 ° C., and a ductile fracture surface ratio at −20 ° C. in a DWTT test of 85% or more .

  According to the present invention, a high-strength line pipe steel plate having a high tensile strength of 900 MPa or more and excellent in strength-toughness balance, DWTT characteristics, and CTOD characteristics is manufactured with high efficiency and at low cost. Can be achieved, and the industrial effect is remarkable.

The present invention defines the composition of steel and production conditions. % In the composition is% by mass.
[Ingredient composition]
C: 0.04 to 0.12%
C is an element that increases the strength of steel, and in order to obtain a desired high strength, it needs to contain 0.04% or more. On the other hand, if the content exceeds 0.12%, the weldability is deteriorated, weld cracking is likely to occur, and the base metal toughness and the HAZ toughness are lowered.

  For this reason, C is limited to the range of 0.04 to 0.12%. In addition, Preferably it is 0.04 to 0.06%.

Si: ≦ 0.50%
Si is an element that acts as a deoxidizing material and further increases the strength of the steel material by solid solution strengthening. However, the content exceeding 0.50% significantly deteriorates the HAZ toughness. Therefore, Si is set to ≦ 0.50%. In addition, Preferably, it is 0.05 to 0.20%.

Mn: 1.80 to 2.50%
Mn is an element that has the effect of improving the hardenability of the steel and improving the toughness. In the present invention, Mn is required to be contained in an amount of 1.80% or more. May deteriorate. For this reason, in this invention, Mn is limited to 1.80 to 2.50% of range. In addition, Preferably, it is 1.80%-2.20%.

P: 0.010% or less P is an element that increases the strength by solid solution strengthening. However, in order to deteriorate toughness and weldability, it is preferable to reduce as much as possible in the present invention. Is acceptable. For this reason, P is limited to 0.010% or less.

S: 0.0020% or less S is present as a sulfide in steel and has an effect of reducing ductility. For this reason, it is desirable to reduce S as much as possible, but it allows up to 0.0020%.

Al: 0.01 to 0.08%
Al acts as a deoxidizer during steelmaking, and in the present invention, it needs to be contained in an amount of 0.01% or more. However, if it exceeds 0.08%, the toughness is reduced. For this reason, Al is limited to the range of 0.01 to 0.08%. In addition, Preferably, it is 0.01 to 0.05%.

Cu: 0.01-0.8%, Cr: 0.01-0.8%, Mo: 0.01-0.8%
Cu, Cr and Mo all act as hardenability improving elements, and if less than 0.01%, the effect cannot be obtained. These are used for replacement of a large amount of Mn, and similarly contribute to increase the strength of the base material and HAZ by obtaining a low-temperature transformation structure, but they are expensive elements and each is 0.8% or more. Even if added, the effect of increasing the strength is saturated, so the upper limit is made 0.8%.

Ni: 0.1 to 1.0%
Ni is also a useful element because it acts as a hardenability improving element and does not cause toughness deterioration when added. In order to obtain this effect, addition of 0.1% or more is necessary, but since it is an expensive element, the upper limit is made 1.0%.

Nb: 0.01 to 0.08%, V: 0.01 to 0.10%
Nb and V are elements necessary for obtaining a required HAZ strength by forming carbides and preventing temper softening of HAZ that is subjected to two or more thermal cycles. In order to obtain these effects, addition of 0.01% or more is necessary.

  Nb also has the effect of expanding the austenite non-recrystallized region during hot rolling. In particular, Nb needs to be added in an amount of 0.01% or more in order to obtain an unrecrystallized region up to 950 ° C. On the other hand, if added over 0.08%, the toughness of the HAZ is remarkably impaired, so the upper limit was made 0.08%. Similarly, if V is added in excess of 0.10%, the toughness of the HAZ is remarkably impaired, so the upper limit is made 0.10%.

Ti: 0.005-0.025%
Ti forms nitrides and is effective in reducing the amount of solute N in the steel. In addition, the precipitated TiN prevents the austenite grains from coarsening by the pinning effect and contributes to the improvement of the toughness of the base material and HAZ.

  Addition of 0.005% or more is necessary to obtain the required pinning effect, but if added over 0.025%, carbides are formed, and the toughness deteriorates significantly due to precipitation hardening. The upper limit is 0.025%.

B: 0.0005 to 0.0030%
B segregates at the austenite grain boundaries and suppresses ferrite transformation, thereby contributing particularly to the prevention of HAZ strength reduction. In order to obtain this effect, addition of 0.0005% or more is required, but even if added over 0.0030%, the effect is saturated, so the upper limit is made 0.0030%.

Ca: ≦ 0.01%
Ca is an element effective in controlling the form of sulfide in steel, and when added, suppresses the generation of MnS harmful to toughness. However, if added over 0.01%, a CaO-CaS cluster is formed and the toughness is deteriorated, so the upper limit is made 0.01%.

REM: ≦ 0.02%
REM is also an effective element for controlling the form of sulfide in steel, and when added, it suppresses the generation of MnS harmful to toughness. However, since it is an expensive element and the effect is saturated even if added over 0.02%, the upper limit is made 0.02%.

N: 0.001 to 0.006%
N is usually present as an inevitable impurity in steel, but TiN that suppresses austenite coarsening is formed by adding Ti as described above. In order to obtain the required pinning effect, 0.001% or more must be present in the steel. However, if it exceeds 0.006%, it was heated to 1450 ° C or more in the weld zone, particularly in the vicinity of the melting line. When TiN decomposes in HAZ, the upper limit is made 0.006% because the solute N has a bad influence.

[Production conditions]
The molten steel having the above composition is melted by a normal melting means such as a converter or an electric furnace, and is made into a steel material such as a slab by a normal casting method such as a continuous casting method or an ingot-bundling method. Is preferred. The melting method and the casting method are not limited to the methods described above.

  The steel material is heated to a temperature at which it becomes an austenite single phase structure. The heating temperature of the steel material is preferably 1000 to 1200 ° C. in order to austenite the steel material. If the heating temperature of the steel material is less than 1000 ° C., the hot deformation resistance is too high, so that the rolling reduction per time cannot be taken high, and the productivity is lowered. Further, when a precipitate-forming element such as V or Nb is contained, these elements are not sufficiently dissolved in austenite, and it is difficult to sufficiently exhibit the effects of these elements.

  On the other hand, when the heating temperature exceeds 1200 ° C., the crystal grains become coarse, and the amount of scale loss and the frequency of furnace repairs increase. For this reason, the heating temperature of the steel material was limited to a range of 1000 to 1200 ° C.

The heated steel material is subjected to hot rolling with a cumulative reduction amount ≧ 70% in a temperature range of 950 ° C. or less and a rolling end temperature in the temperature range of the Ar ₃ transformation point or higher. When the characteristics are desired, hot rolling is performed in which the cumulative reduction rate in the austenite non-recrystallization temperature range is 50% or more and the rolling end temperature is a temperature in the temperature range equal to or higher than the Ar ₃ transformation point.

  If the cumulative rolling reduction in the austenite non-recrystallization temperature region is less than 50%, a fine structure necessary for excellent CTOD characteristics cannot be obtained.

If the rolling end temperature is lower than the Ar ₃ transformation point, ferrite precipitates during rolling, and a desired structure cannot be obtained even if quenching is performed thereafter, and a desired strength cannot be ensured.

After the hot rolling is completed, the steel sheet is quenched and cooled from a temperature range not lower than the Ar ₃ transformation point. When the quenching cooling start temperature is less than the Ar ₃ transformation point, the structure at the quenching cooling start is not an austenite single phase, and the transformation to ferrite or the like has started partially. The amount of sites is small and the desired strength cannot be ensured.

The quenching cooling rate is a cooling rate equal to or higher than the martensite formation critical cooling rate Vcrm. In the present invention, the martensite formation critical cooling rate Vcrm indicates a cooling rate defined by the following equation (1).
logVcrm = 2.94-0.75β
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo) (1)
(Where Vcrm: critical martensite cooling rate (° C./s))
“Martensite formation critical cooling rate Vcrm” means a cooling rate that contains a martensite structure in a fraction of 90% or more of the entire structure.

  In the present invention, each position in the sheet thickness direction is obtained by performing a quenching process for cooling to a quenching and cooling stop temperature in a temperature range of 300 ° C. or more at a martensite transformation start temperature Ms or less at a cooling rate of martensite generation critical cooling rate Vcrm or more. In part, martensite is first generated.

  Martensite is partially generated, and strain is generated at the interface between the generated martensite and untransformed austenite, utilizing expansion during martensitic transformation.

  Strain energy makes it easier for untransformed austenite to transform to lower bainite, and the lower bainite phase can be produced in a finer and larger amount than in the prior art.

  If the quenching cooling rate is less than the martensite formation critical cooling rate Vcrm, the amount of coarse bainite generated before the martensite transformation increases, and the strain generation due to the martensite transformation becomes insufficient, resulting in the expected effect. You can't expect.

  In addition, when the quenching and cooling stop temperature exceeds the Ms point, the strain generation effect due to the formation of martensite cannot be expected, the transformation to the lower bainite phase is insufficiently promoted, and during the subsequent isothermal holding or air cooling. The amount of island martensite harmful to toughness is increased.

  On the other hand, when the quenching and cooling stop temperature is less than 300 ° C., the diffusion of C becomes insufficient, and carbide effective for crack propagation resistance does not precipitate inside the bainitic ferrite.

  Accordingly, the quenching and cooling stop temperature is set to a temperature in the temperature range of 300 ° C. or less below the Ms point. In addition, Preferably, it is a temperature range below 350 degreeC below Ms point.

  Next, after the cooling is stopped at the quenching cooling stop temperature in the above-described range, the steel temperature is maintained within the cooling stop temperature ± 50 ° C. for 60 to 300 seconds, and then cooled to room temperature.

  By maintaining the quenching and cooling stop temperature within ± 50 ° C for 60 to 300 s, martensite is self-annealed, while the transformation of untransformed austenite to lower bainite is promoted, and a mixed structure of tempered martensite and lower bainite is obtained. It is done.

  If the maintained time exceeds 300 s, the matrix structure becomes coarse and the strength decreases. On the other hand, if the maintenance is less than 60S, the transformation to lower bainite is not sufficient, and the target strength and toughness cannot be obtained. For this reason, the maintenance time in the temperature range is limited to a range of 60 to 300 s.

In the present invention, the temperature of the thick steel plate is maintained for 60 to 300 s within the cooling stop temperature ± 50 ° C., and then immediately rises from the temperature to a temperature range of 450 ° C. or more and the Ac ₁ transformation point by ₁ ° C./s or more. By tempering by rapid heating at a high temperature rate, martensite is self-annealed, while transformation of untransformed austenite to lower bainite is promoted, and a mixed structure of tempered martensite and lower bainite can be obtained.

This makes it possible to improve toughness with almost no deterioration in strength. When the heating temperature is lower than 450 ° C., the effect of improving toughness can not be almost obtained, since when the Ac ₁ transformation point or more temperature reduction in strength occurs, the heating temperature is less than Ac ₁ transformation point 450 ° C. or higher.

  Further, when the rate of temperature rise is less than 1 ° C./s, the toughness is improved, but the strength is significantly reduced. For this reason, the heating rate is set to 1 ° C./s or more.

  The thick steel plate obtained under the manufacturing conditions described above has the above-described composition, and has a mixed structure of tempered martensite and lower bainite regardless of the position in the plate thickness direction. The structural fraction of tempered martensite and lower bainite is 90% or more by volume. In addition, it is preferable that the structure fraction of a lower bainite becomes 50% or more of the whole by volume ratio.

  Further, as phases other than tempered martensite and lower bainite, it is allowable to mix upper bainite and ferrite having a volume ratio of 10% or less. Here, “tempered martensite” refers to martensite in which carbides are precipitated or spheroidized.

  “Lower bainite” includes tempered lower bainite in which carbides are precipitated or spheroidized.

  In actual operation, the temperature of the steel sheet is generally controlled by the surface temperature of the steel sheet. In general, the average temperature of the entire steel sheet is calculated in real time, and temperature control and speed control are performed based on this average temperature. In the present invention, “temperature” means the average temperature of the entire steel sheet, “cooling rate” means the average cooling rate of the entire steel sheet, and “heating rate” means the average heating rate of the entire steel sheet.

  The present invention uses a facility in which a high-frequency induction heating device is disposed as a heating device in the case of maintaining a constant temperature after quenching or reheating on a transport line downstream of a cooling device for quenching. The heat history after quenching stop, in particular, quenching is stopped at a temperature below the Ms point and maintained for a short time in the temperature range, which is very effective.

There are no particular restrictions on the steel making method, but from the economical point of view, it is desirable to cast steel pieces by the converter method and by continuous casting.
In the present _invention, Ar 3, each transformation point Ac ₁ is the following equation based on the content of each element in each steel material (steel plate) (4), (5)
Ar ₃ = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo (4)
_{Ac 1 = 751-26.6C + 17.6Si-11.6Mn} -169Al-23Cu-23Ni + 24.1Cr + 22.5Mo-39.7V + 233Nb-5.7Ti-895B
... (5)
The value obtained by calculating using is used.

  Steel plates 1-1 to 9 were produced under the hot rolling / accelerated cooling / online heating conditions shown in Table 2 using steel having the chemical composition shown in Table 1. In addition, the steel plate temperature and cooling rate in Table 2 indicate average temperature and average cooling rate.

  From the steel plate obtained under the manufacturing conditions shown in Table 2, a full thickness tensile test piece conforming to API-5L, a DWTT test piece, and a V-notch Charpy impact test piece of JIS Z2202 from the center position of the plate thickness were collected and the steel plate was pulled. Tests, DWTT tests and Charpy impact tests were conducted to evaluate strength and toughness.

  In addition, a specimen for tissue observation was collected, and the structure was observed at a position in the plate thickness direction 1/2 with a scanning electron microscope and a transmission electron microscope to identify the tissue and obtain the tissue fraction of each tissue.

  In addition, tempered martensite and lower bainite were discriminated by the precipitation form of carbides. The structure fraction of each structure is determined by measuring the average austenite (γ) particle size by a line segment method using a scanning electron microscope, and randomly selecting 10 particles having the average γ particle size. The area of each structure was determined as the cross-sectional area ratio, and the average value of the ten cross-sectional area ratios was taken as the structure fraction at each position of the steel sheet.

  Table 3 summarizes the evaluation results of the strength and toughness of the base metal.

  Steel sheet chemical composition within the scope of the present invention, the present invention example of rolling conditions shows a base metal strength exceeding 900 MPa, and a high toughness exceeding 200 J in a Charpy impact test at -30 ° C, and in a DWTT test at -20 ° C A ductile fracture surface ratio of 85% or more was obtained.

  On the other hand, Comparative Example No. in which the cooling start temperature was below the lower limit of the present invention. In 1-3, the strength decreased because a part of the structure was first transformed into ferrite.

  Comparative Example No. in which the cooling stop temperature after rolling exceeded the upper limit of the present invention. In 2-3, martensite transformation did not occur, and a bainite single-phase structure was formed. Further, as a result of coarsening of the bainite lower structure due to cooling stop at a higher temperature, the strength decreased.

  Comparative Example No. in which the cooling stop temperature after rolling was lower than the lower limit of the present invention. Since 3-2 became a tempered martensite main structure instead of the lower bainite main structure, the strength was high, but the Charpy absorbed energy and DWTT characteristics were lowered.

Comparative Example No. in which the cooling rate after rolling was lower than the lower limit of the present invention. In 4-3, the strength was significantly reduced. Comparative Example No. in which the online heating temperature after cooling stopped exceeded the upper limit of the present invention. As for 4-4, as a result of exceeding the Ac ₁ transformation point of the steel sheet, α-γ reverse transformation occurred and MA was formed, but the amount was not sufficient and the strength was lowered.

  Comparative Example No. in which the holding time at the cooling stop temperature ± 50 ° C. exceeded the upper limit of the present invention. In 5-3, the base material strength and the DWTT characteristics were lowered. Comparative example No. in which the rate of temperature increase during online heating after cooling stopped was below the lower limit of the present invention. In 6-2, the base material strength showed a high value, but the Charpy absorbed energy and the DWTT characteristic were lowered.

  Comparative Example No. B added amount below the lower limit of the present invention. In No. 7, the strength decreased as a result of insufficient hardenability.

  Comparative Example No. in which the Mn addition amount of the steel sheet was below the lower limit of the present invention. Even at 8, the strength decreased. On the other hand, Comparative Example No. in which the C addition amount of the steel sheet exceeded the upper limit of the present invention. Although No. 9 showed high strength, strength and YR became too high, and Charpy absorbed energy and DWTT characteristics were lowered.

  As described above, it was confirmed that the comparative examples whose component composition and / or production conditions were outside the scope of the present invention were inferior to the characteristics of the present invention.

  Moreover, using the steel type D in Table 1, steel sheets 4-1, 4-2, 4-5, 4-6 were produced by changing the cumulative rolling reduction in the untransformed austenite region in the hot rolling, and described above. In addition to the test items, a three-point bending CTOD test piece of B (plate thickness) × 2B size in accordance with BS7748 was collected and a CTOD test was performed.

  Table 4 shows the manufacturing conditions, and Table 5 shows the test results. No. 3 which is an example of the present invention according to claim 3. In 4-1 and 4-2, a critical opening displacement amount of 0.15 mm or more at a test temperature of −20 ° C. was obtained in the CTOD test.

  On the other hand, Comparative Example No. in which the cumulative rolling reduction in the austenite non-recrystallized region was lower than the lower limit of the present invention described in claim 3. In 4-5 and 4-6, austenite grains are not sufficiently refined, and the CTOD characteristics are inferior to those of the examples of the present invention.

Claims (4)

In mass%, C: 0.04 to 0.12%, Si: ≦ 0.50%, Mn: 1.80 to 2.5
0%, P ≦ 0.010%, S ≦ 0.002%, Al: 0.01 to 0.08%, Cu: 0.01 to 0.8%, Ni: 0.1 to 1.0%, Cr: 0.01-0.8%, Mo: 0.01-0.8%, Nb: 0.01-0.08%, V: 0.01-0.10%, Ti: 0.005- 0.025%, B: 0.0005 to 0.0030%, Ca: ≦ 0.01%, REM: ≦ 0.02%, N: 0.001 to 0.006%, the balance Fe and inevitable After heating the steel consisting of mechanical impurities to 1000-1200 ° C., hot rolling is performed at a rolling end temperature of Ar ₃ transformation point or higher and a cumulative reduction of 50% or higher in the austenite non-recrystallization temperature range. ₃ transformation point or higher temperature range, in martensite critical cooling rate Vcrm or more cooling rate defined by the following formula (1) After cooling to the quenching and cooling stop temperature in the temperature range of 300 ° C or higher below the martensite transformation start temperature (Ms point) defined by the following formula (2), the steel is heated online for 60s to 300s after the cooling stop. The tensile strength is 900 MPa or more, the Charpy impact test at −30 ° C. exceeds 200 J, and the DWTT test is −20. A method for producing a steel sheet for a high-tensile line pipe having a characteristic of a ductile fracture surface ratio at 85 ° C. of 85% or more .
logVcrm = 2.94-0.75β
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo)
... (1)
Here, Vcrm: Martensite formation critical cooling rate (° C./s)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
... (2) In mass%, C: 0.04 to 0.12%, Si: ≦ 0.50%, Mn: 1.80 to 2.5
0%, P ≦ 0.010%, S ≦ 0.002%, Al: 0.01 to 0.08%, Cu: 0.00.
01-0.8%, Ni: 0.1-1.0%, Cr: 0.01-0.8%, Mo: 0.01
-0.8%, Nb: 0.01-0.08%, V: 0.01-0.10%, Ti: 0.00
5 to 0.025%, B: 0.0005 to 0.0030%, Ca: ≦ 0.01%, REM:
≦ 0.02%, N: 0.001 to 0.006%, and the steel consisting of the remainder Fe and inevitable impurities is heated to 1000 to 1200 ° C., and the rolling end temperature is at least the Ar ₃ transformation point . Hot rolling with a cumulative reduction ratio of 50% or more in the austenite non-recrystallization temperature range, and then a martensite formation critical cooling rate Vcrm defined by the following formula (1) from a temperature range of Ar ₃ transformation point or higher After cooling to the quenching cooling stop temperature in the temperature range of 300 ° C. or higher, which is equal to or lower than the martensite transformation start temperature (Ms point) defined by the following formula (2), for 60 s to 300 s after the cooling stop, the online heating, the temperature of the steel is maintained within the cooling stop temperature ± 50 ° C., then, to rapid heating immediately from the temperature to 450 ° C. or higher Ac ₁ transformation point of the temperature range at 1 ° C. / s or more heating rate Features and tensile strength than 900MPa that, in more than 200J In Charpy impact test at -30 ° C., the ductility fracture rate at -20 ° C. In DWTT test for high tension line pipe having a 85% or more properties A method of manufacturing a steel sheet.
logVcrm = 2.94-0.75β
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo)
... (1)
Here, Vcrm: Martensite formation critical cooling rate (° C./s)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
... (2) Heating after water cooling is performed by a high frequency induction heating device installed on the same line as the cooling device on the downstream side of the cooling device, and the tensile strength according to claim 1 or 2 is 900 MPa or more, A method for producing a steel sheet for a high-strength line pipe having a characteristic that the Charpy impact test at −30 ° C. exceeds 200 J and the ductile fracture surface ratio at −20 ° C. is 85% or more in the DWTT test . It is manufactured by the manufacturing method according to any one of claims 1 to 3 , wherein the microstructure is a mixed structure of tempered martensite and a lower bainite having a volume ratio of 90% or more, and the lower bainite is at least a volume ratio of 50% or more. A steel sheet for high-strength line pipes , having a tensile strength of 900 MPa or more, a characteristic of exceeding 200 J in a Charpy impact test at −30 ° C. and a ductile fracture surface ratio at −20 ° C. in a DWTT test of 85% or more .