JP5672658B2 - High strength steel plate for line pipes with excellent HIC resistance and weld heat affected zone toughness and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength steel sheet for line pipes excellent in hydrogen-induced crack resistance (HIC resistance) and weld heat affected zone toughness, and a method for producing the same.

  Line pipes used to transport crude oil and natural gas containing hydrogen sulfide have strength, toughness and weldability, as well as hydrogen-induced crack resistance (HIC resistance) and stress corrosion crack resistance (SCC resistance). So-called sour resistance is required. In hydrogen induced cracking (HIC) of steel, hydrogen ions from the corrosion reaction are adsorbed on the surface of the steel, penetrate into the steel as atomic hydrogen, and include non-metallic inclusions such as MnS in the steel and hard second phase structure. It is said that it diffuses and accumulates around, and cracks occur due to its internal pressure.

  In order to prevent such hydrogen-induced cracking, in Patent Document 1, by adding an appropriate amount of Ca or Ce to the amount of S, the formation of acicular MnS is suppressed, and the finely dispersed spherical shape with a small stress concentration A method for producing steel for line pipes having excellent HIC resistance is disclosed in which the shape of the inclusions is changed to suppress the generation and propagation of cracks. In Patent Documents 2 and 3, the central segregation part is reduced by reducing elements that have a high segregation tendency (C, Mn, P, etc.), soaking in the slab heating stage, and accelerated cooling during transformation during cooling. Steel having excellent HIC resistance is disclosed that suppresses the formation of hardened structures such as island martensite, which is the starting point of cracks, and martensite, bainite, which is the propagation path of cracks.

  In addition, as a high-strength steel excellent in HIC resistance that does not contain block-like bainite or martensite whose microstructure is highly susceptible to cracking, Patent Document 4 describes the HIC resistance of APIX 80 grade, which is a ferrite-bainite two-phase structure. A high-strength steel material excellent in the above is disclosed. In Patent Documents 5 and 6, SCC (SSCC) resistance and HIC resistance are improved by making the microstructure a ferrite single phase structure, and the precipitation strengthening of carbide obtained by adding a large amount of Mo or Ti is used. A high strength steel is disclosed.

  On the other hand, the demand for steel sheets having higher strength and higher toughness is increasing from the viewpoint of increasing the size of welded steel structures and reducing costs. Usually, a high-strength and high-toughness steel sheet is manufactured by a so-called TMCP method using quenching and tempering treatment or controlled rolling / controlled cooling. However, the quenching and tempering treatment requires time and labor and is expensive to manufacture. Moreover, when increasing the strength of steel materials using the TMCP method, it is necessary to add a large amount of alloy elements to the steel materials, which raises costs due to the addition of alloy elements and deteriorates the toughness of the weld heat affected zone. .

Japanese Patent Laid-Open No. 54-110119 JP 61-60866 A JP-A-61-165207 JP 7-216500 A Japanese Patent Laid-Open No. 61-227129 JP-A-7-70697

  The methods for improving the HIC resistance described in Patent Documents 1 to 3 are all about the center segregation part. Since high strength steel plates exceeding the APIX 70 grade are often manufactured by accelerated cooling or direct quenching, the steel plate surface portion having a high cooling rate is hardened compared to the inside, and hydrogen-induced cracks are generated from the vicinity of the surface. In addition, the microstructure of these high-strength steel sheets obtained by accelerated cooling is a structure with relatively high susceptibility to cracking of bainite or acicular ferrite not only on the surface but also inside, so that the HIC resistance of high-strength steel exceeding APIX 70 grade is high. In addition to the need for further strict segregation suppression, it is necessary to take measures against HIC starting from sulfides or oxide inclusions. Therefore, when considering the HIC resistance of these high-strength steel sheets, in addition to the component design and microstructure control that enables strict segregation suppression while ensuring the strength, the HIC of the surface part of the steel sheet, It is necessary to take measures against HIC starting from sulfides and oxide inclusions.

The bainite structure of high-strength steel described in Patent Document 4 is a structure that is not as high as block bainite or martensite but is relatively high in cracking sensitivity. S and Mn amounts are severely limited, and Ca treatment is essential. Since it is necessary to improve the HIC property, the manufacturing cost is high. Moreover, it is difficult to stably obtain a ferrite-bainite two-phase structure using the rolling / cooling method described in Patent Document 4.
On the other hand, since the ferrite phase of the steels described in Patent Documents 5 and 6 is a structure rich in ductility and has extremely low cracking susceptibility, the HIC resistance is greatly improved compared to steels having a bainite structure or an acicular ferrite structure. . However, since the strength of the ferrite single phase is low, Patent Document 5 increases the strength by adding a large amount of C and Mo and precipitates a large amount of carbide, while Patent Document 6 increases the strength of Ti-added steel at a specific temperature. It is rolled up on a steel strip and strengthened using TiC precipitation strengthening.

  However, in order to obtain the ferrite structure in which the Mo carbides described in Patent Document 5 are dispersed, it is necessary to perform cold working after quenching and tempering, and then perform tempering again, which increases the manufacturing cost. Moreover, since the particle size of Mo carbide is as large as about 0.1 μm and the effect of increasing the strength is low, it is necessary to obtain a predetermined strength by increasing the content of C and Mo and increasing the amount of carbide. In addition, TiC used in the high-strength steel described in Patent Document 6 is finer than Mo carbide and is an effective carbide for precipitation strengthening, but is easily coarsened under the influence of temperature during precipitation. In Patent Document 6, no countermeasure is taken against such coarsening of precipitates. For this reason, precipitation strengthening is not sufficient, and a large amount of Ti is required. Adding a large amount of alloy elements not only increases the material cost but also degrades the weld heat-affected zone toughness, which is undesirable when high toughness is required.

Accordingly, an object of the present invention is to solve such problems of the prior art, and is a high-strength steel sheet for line pipes having a strength of APIX 70 grade or higher (tensile strength of 580 MPa or higher). A line pipe that has excellent HIC resistance against HIC generated from inclusions, has excellent weld heat affected zone toughness, and can be manufactured at low cost without adding a large amount of alloying elements. It is to provide a high strength steel sheet for use.
Moreover, the other object of this invention is to provide the manufacturing method which can manufacture stably the high strength steel plate for line pipes which has such the outstanding performance.

The features of the present invention for solving such problems are as follows.
[1] By mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 1.8%, P: 0.01% or less, S: 0.002% or less, Ca: 0.0005 to 0.005%, Nb: 0.05 to 0.15%, Al: 0.01 to 0.08%, and further, V: 0.005 to 0.15%, Ti: 0.005 to 0.04% of one or two kinds, the balance being Fe and inevitable impurities, and the amount of C in atomic% and the sum of Nb, V and Ti C / (Nb + V + Ti) which is the ratio of the amount is 1.0 to 5.0, CP value (mass%) represented by the following formula (1) is 0.98 or less, P represented by the following formula (2) Having a component composition with a _CM value (mass%) of 0.15 or less,
The metal structure is a substantial two-phase structure in which the sum of the ferrite phase and the bainite phase is 95% or more by volume fraction, and carbides containing one or two of Nb, V, and Ti are dispersed and precipitated. A high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness, characterized by having a tensile strength of 580 MPa or more.
CP = 4.46C (%) + 2.37Mn (%) / 6+ {1.74Cu (%) + 1.7Ni (%)} / 15+ {1.18Cr (%) + 1.95Mo (%) + 1.74V (% )} / 5 + 22.36P (%) (1)
P _CM = C (%) + Si (%) / 30 + Mn (%) / 20 + Cu (%) / 20 + Ni (%) / 60 + Cr (%) / 20 + Mo (%) / 15 + V (%) / 10 + B (%) * 5 (5) 2)
However, in the formulas (1) and (2), the element not added is 0.

[2] In the high-strength steel sheet for line pipes of the above [1], further, by mass, Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less, Mo: 0.00. A high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness, characterized by containing one or more of 50% or less and B: 0.005% or less.
[3] A method for producing a high-strength steel sheet for line pipes according to [1] or [2 ] above,
The steel having the component composition described in [1] or [2] above is heated to a temperature of 1000 to 1300 ° C., hot-rolled at a rolling end temperature of Ar ₃ temperature or higher, and then cooled at 5 ° C./sec or higher. Accelerating cooling to 300 to 600 ° C. at a speed, and immediately after that, at a temperature rising rate of 0.5 ° C./sec or more, reheating to 550 to 700 ° C. above the cooling stop temperature, A method for producing high-strength steel sheets for line pipes with excellent HIC resistance and weld heat-affected zone toughness.
[4] A high-strength steel pipe for a line pipe excellent in HIC resistance and weld heat affected zone toughness, characterized by being manufactured using the steel sheet described in [1] or [2] above.

The high-strength steel sheet for line pipes of the present invention has excellent HIC resistance and weld heat affected zone toughness while having high strength, and can be manufactured at a low cost without adding a large amount of alloying elements.
Moreover, according to the manufacturing method of this invention, the high strength steel plate for line pipes which has the above high intensity | strength, the outstanding HIC resistance characteristic, and the weld heat affected zone toughness can be manufactured stably.
The high-strength steel pipe for line pipe of the present invention has high strength and excellent HIC resistance and weld heat affected zone toughness, and is therefore suitable for transporting crude oil and natural gas containing hydrogen sulfide.

The graph which shows the outline of the heat history for the metal structure control in the manufacturing method of this invention Explanatory drawing which shows an example of the manufacturing line for enforcing the manufacturing method of this invention

  The present inventors diligently studied the components and microstructure of the steel material and the manufacturing method of the steel sheet in order to achieve both improvement in the HIC resistance and high strength of the high-strength steel sheet and further improve the weld heat affected zone toughness. . As a result, in order to achieve both high strength and anti-HIC properties, the CP value in consideration of segregation is optimized so that the segregation is suppressed and the crack resistance is low, while maintaining high strength. It is most effective to have a ferrite + bainite two-phase structure with a small strength difference between the ferrite phase and the bainite phase. By adopting a manufacturing process of accelerated cooling after controlled rolling and subsequent reheating, Ti, Nb, It was found that the ferrite phase, which is a soft phase, and softening of the bainite phase, which is a hard phase, are caused by fine precipitates including V and the like, and a ferrite + bainite two-phase structure with a small strength difference can be obtained. And in addition to strengthening by bainite transformation during accelerated cooling, by optimizing the amount of Ti, Nb, V added to C, precipitation strengthening due to fine carbides that precipitate during reheating can be maximized, It was found that high strength can be achieved even in low-component steels with few alloying elements. In particular, it has been found that by effectively utilizing Nb, the effects of transformation strengthening and precipitation strengthening can be increased and the structure can be refined.

  A high-strength steel sheet having a two-phase structure of a ferrite phase and a bainite phase in which precipitates containing Ti, Nb, V and the like as described above are dispersed is a steel sheet having a bainite or acicular ferrite structure obtained by conventional accelerated cooling or the like. Thus, there is no increase in hardness at the surface layer portion, so that HIC from the surface layer portion does not occur. Furthermore, the two-phase structure of the ferrite phase and the bainite phase having a small strength difference has extremely high resistance to cracking, and it is possible to suppress HIC from the steel plate center and inclusions. Moreover, the crack from a center segregation part can be suppressed by managing strictly the amount of alloy components with a segregation tendency, and regulating by the CP value. Furthermore, in order to make maximum use of precipitation strengthening in addition to transformation strengthening, it is not necessary to add a large amount of alloy elements, and high strength can be achieved without impairing the toughness of the weld heat affected zone.

Hereinafter, the high-strength steel sheet of the present invention will be described in detail. First, the structure of the high-strength steel sheet of the present invention will be described.
The metal structure of the high-strength steel sheet of the present invention is a substantially two-phase structure in which the total of the ferrite phase and the bainite phase is 95% or more in volume fraction. Since the ferrite phase is rich in ductility and has low cracking susceptibility, high HIC resistance can be realized. The ferrite of the high-strength steel sheet of the present invention is a fine granular ferrite or bainitic ferrite in which untransformed austenite remaining after accelerated cooling is transformed into ferrite, and has a smooth and clear grain boundary. Excellent strength and toughness compared to Nalferrite. The bainite phase has excellent strength and toughness due to transformation strengthening. The two-phase structure of ferrite and bainite is generally a mixed structure of a soft ferrite phase and a hard bainite phase, and in a steel material having such a structure, hydrogen is likely to accumulate at the interface between the ferrite phase and the bainite phase. In addition, since the interface serves as a crack propagation path, the HIC resistance is inferior. On the other hand, in the present invention, by adjusting the strength of the ferrite phase and the bainite phase to reduce the difference in strength between them, it is possible to achieve both HIC resistance and high strength.

When two or more different metal structures such as martensite, pearlite, retained austenite, and island martensite (MA) are mixed in the ferrite + bainite two-phase structure, hydrogen accumulation and stress concentration at the heterophase interface Since it becomes easy to generate HIC, the smaller the structure other than the ferrite phase and the bainite phase, the better. However, when the volume fraction of the structure other than the ferrite phase and the bainite phase is sufficiently low, the influence thereof can be ignored. Specifically, the sum of the metal structures other than the ferrite phase and the bainite phase (one or more of martensite, pearlite, retained austenite, island martensite (MA), etc.) is less than 5% in volume fraction. If there is, there is no big influence. Therefore, the metal structure of the steel sheet of the present invention may be a substantial two-phase structure in which the total of the ferrite phase and the bainite phase is 95% or more in terms of volume fraction. In particular, from the viewpoint of HIC resistance, the volume fraction of island martensite (MA) is more preferably 3% or less.
The volume fraction of bainite is not particularly defined, but is preferably 10% or more from the viewpoint of securing the toughness of the base material, and preferably 80% or less from the viewpoint of HIC resistance, and a more preferable volume fraction is 20 to 60%. It is.

Next, in the high-strength steel sheet of the present invention, the precipitate that is dispersed and precipitated in the ferrite phase will be described.
The high-strength steel sheet of the present invention utilizes precipitation strengthening due to a composite carbide containing Nb in the ferrite phase and one or two of V and Ti (V and / or Ti). Fine precipitation of composite carbide containing Nb and one or two of V and Ti in the steel strengthens the ferrite phase and reduces the difference in strength between the ferrite phase and the bainite phase. Characteristics can be obtained. Since this precipitate is extremely fine, it has no influence on the HIC resistance.
Nb, V, and Ti are elements that form carbides in steel, and strengthening of steel by precipitation of individual carbides has been conventionally performed, but conventionally, by cooling process and isothermal holding after hot rolling. Utilizing precipitation from austenite or precipitation from supersaturated ferrite, or quenching after rolling to make the structure martensite or bainite, then tempering in a heating furnace to convert carbide into martensite or bainite The method of making it precipitate was taken.

On the other hand, the present invention deposits carbide using rapid heating using an induction heating furnace or the like in the reheating process from the bainite transformation region. According to such a method of the present invention, the coarsening of the carbide is suppressed by heating in a rapid time and a very fine carbide is precipitated, so that a greater strength improvement effect is obtained as compared with the normal method. It is characteristic that Such an unprecedented strength improvement effect is extremely fine with a particle size of less than 10 nm because the composite carbide containing one or two of Nb, V, and Ti is stable and the growth rate is slow due to rapid heating. It is because it is obtained as a simple precipitate. Such precipitates having a particle size of less than 10 nm are preferably deposited at 2 × 10 ³ pieces / μm ³ or more in order to obtain a high-strength steel sheet having a tensile strength of 580 MPa or more (APIX 70 grade or more). The form of precipitation may be random or in line, and there is no particular limitation. Moreover, although this fine carbide precipitates mainly in a ferrite phase, it may precipitate also from a bainite phase depending on a chemical component and manufacturing conditions. In this case, further strengthening is possible, but if the strength difference between the ferrite phase and the bainite phase is HV70 or less, the HIC resistance is not affected.
The composite carbide containing one or two of Nb, V, and Ti, which is a precipitate that disperses and precipitates in the steel sheet of the present invention, is applied to the steel having the component composition described below. Can be obtained.

Next, the component composition of the high-strength steel sheet of the present invention will be described. In the following description, all units indicated by% are mass%.
C: Set to 0.02 to 0.08%. C is an element that contributes to precipitation strengthening as a carbide, but if it is less than 0.02%, sufficient strength cannot be secured, while if it exceeds 0.08%, toughness is deteriorated.
Si: 0.01 to 0.5%. Si is added for deoxidation, but if it is less than 0.01%, the deoxidation effect is not sufficient, while if it exceeds 0.5%, toughness and weldability are deteriorated.
Mn: 0.5 to 1.8%. Mn is added for strength and toughness, but if it is less than 0.5%, its effect is not sufficient. On the other hand, if it exceeds 1.8%, weldability and HIC resistance are deteriorated. In particular, from the viewpoint of HIC resistance, the preferable amount of Mn is 0.5 to 1.6%.

P: 0.01% or less. P is an unavoidable impurity element, which deteriorates the weldability and deteriorates the HIC resistance by increasing the hardness of the central segregation part, and the tendency becomes remarkable when it exceeds 0.01%. In particular, from the viewpoint of HIC resistance, the preferable amount of P is 0.008% or less.
S: Set to 0.002% or less. In general, the S content becomes MnS inclusions in the steel and deteriorates the HIC resistance, so the smaller the better, but there is no problem if it is 0.002% or less.
Ca: 0.0005 to 0.005%. Ca is an element effective for improving the HIC resistance by controlling the form of sulfide inclusions. However, if it is less than 0.0005%, the effect is not sufficient. On the other hand, even if it exceeds 0.005%, it is effective. Saturates, but rather deteriorates the anti-HIC properties due to a reduction in the cleanliness of the steel.

Nb: 0.05 to 0.15%. Nb is an important element in the present invention. Nb can effectively increase the strength by utilizing both transformation strengthening and precipitation strengthening, and can refine the structure. Nb expands the non-recrystallized region and is an element effective for strengthening transformation, and therefore increases the effect of transformation strengthening and microstructure refinement by TMCP. Further, Nb improves strength and toughness by increasing transformation strengthening and refining the structure, and forms fine composite carbides together with V and Ti, thereby contributing to an increase in strength. However, if it is less than 0.05%, the effect is not sufficient. On the other hand, if it exceeds 0.15%, the toughness of the weld heat affected zone deteriorates. From the viewpoint of fully utilizing transformation strengthening and precipitation strengthening, a preferable Nb amount is 0.07 to 0.15%. Moreover, in order to set the strength of the steel sheet to 620 MPa or more (APIX 80 grade or more) and suppress the toughness deterioration of the weld heat affected zone, the Nb content is preferably 0.08 to 0.12%.
Al: 0.01 to 0.08%. Al is added as a deoxidizer, but if it is less than 0.01%, there is no effect. On the other hand, if it exceeds 0.08%, the cleanliness of the steel decreases and the toughness deteriorates.

1 type or 2 types of V: 0.005-0.15% and Ti: 0.005-0.04% are contained.
V forms a fine composite carbide together with Nb and Ti and contributes to an increase in strength. However, if it is less than 0.005%, the effect is not sufficient. On the other hand, if it exceeds 0.15%, the toughness of the heat affected zone is deteriorated. For this reason, when adding V, it is set as 0.005 to 0.15%. From the viewpoint of fully utilizing precipitation strengthening and suppressing toughness deterioration of the weld heat affected zone, the preferable V amount is 0.005 to 0.12%.
Ti forms a fine composite carbide together with Nb and V, and greatly contributes to an increase in strength. However, if it is less than 0.005%, the effect is not sufficient. On the other hand, if it exceeds 0.04%, the toughness of the weld heat affected zone deteriorates. For this reason, when adding Ti, it is made into 0.005 to 0.04%. From the viewpoints of fully utilizing precipitation strengthening and suppressing toughness deterioration of the weld heat affected zone, the preferable Ti amount is 0.005 to 0.03%.

C / (Nb + V + Ti), which is a ratio of the amount of C in atomic% and the total amount of Nb, V, and Ti, is set to 1.0 to 5.0. In order to effectively use precipitation strengthening by composite precipitates, the relationship between the amount of C and the amount of carbide forming elements Nb, V, and Ti is important, and these elements are added in an appropriate balance. As a result, a thermally stable and very fine composite carbide can be obtained. When the value of this parameter formula is less than 1.0 or more than 5.0, the content of any element is excessive, and a fine composite carbide having a particle size of less than 10 nm cannot be sufficiently obtained. Cause deterioration of HIC resistance and toughness due to the formation of a hardened structure such as martensite.
In addition, when using content by mass%, the value of (C / 12.01) / (Ti / 47.9 + Nb / 92.91 + V / 50.94) shall be 1.0-5.0.

The CP value (mass%) represented by the following formula (1) is set to 0.98 or less. In the following formula (1), C (%), Mn (%), Cu (%), Ni (%), Cr (%), Mo (%), V (%), and P (%) are the respective elements. The content (% by mass) is 0 for elements not added.
CP = 4.46C (%) + 2.37Mn (%) / 6+ {1.74Cu (%) + 1.7Ni (%)} / 15+ {1.18Cr (%) + 1.95Mo (%) + 1.74V (% )} / 5 + 22.36P (%) (1)
This CP value is an equation designed to estimate the material of the center segregation part from the content of each alloy element. The higher the CP value, the higher the concentration of the center segregation part, and the hardness of the center segregation part. To rise. By setting the CP value to 0.98 or less, it is possible to suppress cracks in the HIC test. Further, the lower the CP value, the lower the hardness of the center segregation part. Therefore, when higher HIC resistance is required, the upper limit is desirably set to 0.95.

Following (2) _{P CM} value of the formula (mass%) 0.15 or less. In the following formula (2), C (%), Si (%), Mn (%), Cu (%), Ni (%), Cr (%), Mo (%), V (%), B (% ) Is the content (% by mass) of each element, and the elements not added are 0.
P _CM = C (%) + Si (%) / 30 + Mn (%) / 20 + Cu (%) / 20 + Ni (%) / 60 + Cr (%) / 20 + Mo (%) / 15 + V (%) / 10 + B (%) * 5 (5) 2)
In the present invention, in order to ensure good weld heat-affected zone toughness, a low alloy component composition with a small alloy element amount of P _CM ≦ 0.15 mass% is adopted.

The above is the basic component composition of the present invention. However, when the strength and toughness of the steel sheet are further improved, Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less, Mo : 0.50% or less, B: 0.005% or less may be included.
Cu is an element effective for improving toughness and increasing strength, but if added excessively, weldability deteriorates, so when added, the upper limit is 0.50%.
Ni is an element effective for improving toughness and increasing strength. However, if added excessively, it is disadvantageous in terms of cost, and weld heat-affected zone toughness deteriorates, so when added, the upper limit is 0.50%. And
Cr, like Mn, is an element effective for obtaining sufficient strength even at low C. However, if added excessively, weldability deteriorates, so when added, the upper limit is 0.50%.
Mo is an element effective for improving toughness and increasing strength, but if added excessively, weldability deteriorates, so when added, the upper limit is 0.50.
B is an element that contributes to an increase in strength and an improvement in HAZ toughness, but if added in excess, weldability deteriorates, so when added, the upper limit is made 0.005%.
The balance other than the above consists of Fe and inevitable impurities. However, the content of other trace elements is not hindered unless the effects of the present invention are impaired.

The production method of the high-strength steel sheet of the present invention is not particularly limited, but the following production method is desirable. In the following description, the temperature is an average temperature of a slab or a steel plate.
In FIG. 1, the outline of the heat history for the metal structure control in the manufacturing method of this invention is shown. In the production method of the present invention, accelerated cooling is performed from the austenite region at an Ar ₃ temperature or higher to the bainite region to form a mixed structure of untransformed austenite and bainite. Transformation occurs and fine precipitates are dispersed and precipitated in the ferrite phase. On the other hand, the bainite phase is tempered to become tempered bainite. By adopting a two-phase structure consisting of a ferrite phase precipitation strengthened by these fine precipitates and a tempered and softened bainite layer, it is possible to achieve both high strength and HIC resistance without adding a large amount of alloying elements. Become.

In the production method of the present invention, the steel (slab) having the above-described component composition is hot-rolled at a predetermined slab heating temperature and rolling end temperature, acceleratedly cooled under predetermined conditions, and then immediately reheated by rapid heating. I do.
Slab heating temperature: 1000-1300 ° C. If the heating temperature is less than 1000 ° C., the solid solution of the carbide is insufficient and the required strength cannot be obtained, while if it exceeds 1300 ° C., the toughness deteriorates.
Hot rolling end temperature: Ar ₃ temperature or higher. The Ar ₃ temperature means a ferrite transformation start temperature during cooling, and can be obtained by the following equation. In the formula, C (%), Mn (%), Cu (%), Cr (%), Ni (%), and Mo (%) are the contents (%) of each element, and the elements not added are 0. To do. When the rolling end temperature is lower than the Ar ₃ temperature, the subsequent ferrite transformation rate is lowered, so that sufficient precipitation of fine precipitates cannot be obtained during ferrite transformation by reheating, and the strength is lowered.
_{Ar 3 = 910-310C (%) -} 80Mn (%) - 20Cu (%) - 15Cr (%) - 55Ni (%) - 80Mo (%)

Immediately after the end of rolling, accelerated cooling to 300 to 600 ° C. is performed at a cooling rate of 5 ° C./sec or more. When the product is allowed to cool or gradually cool after the rolling, the precipitate is easily coarsened because the precipitate is precipitated from the high temperature region, and as a result, sufficient strength cannot be obtained and sufficient transformation strengthening cannot be obtained. . For this reason, in the present invention, rapid cooling (accelerated cooling) is performed to a temperature optimum for precipitation strengthening and transformation strengthening, preventing precipitation from a high temperature region and obtaining the effect of transformation strengthening. There is no particular restriction on the cooling equipment used for this accelerated cooling.
When the cooling rate is less than 5 ° C./sec, the effect of preventing precipitation in a high temperature range is not sufficient, the strength is lowered, and transformation strengthening due to bainite transformation cannot be sufficiently obtained. In addition, ferrite may be generated in a high temperature range during cooling, and precipitates generated during ferrite transformation are easily coarsened in the high temperature range, so that sufficient strength cannot be obtained. In order to sufficiently exhibit the effect of preventing precipitation in a high temperature range and the effect of transformation strengthening by bainite transformation, the cooling rate after the rolling is preferably 10 ° C./sec or more.

  By rapidly cooling to 300 to 600 ° C., which is a bainite transformation region, by accelerated cooling after the end of rolling, a bainite phase is generated, and the driving force for ferrite transformation during reheating is increased. By increasing the driving force, it becomes possible to promote the ferrite transformation in the reheating process and complete the ferrite transformation with a short reheating. When the cooling stop temperature is less than 300 ° C., the island-shaped martensite (MA) is generated even when the single phase structure of bainite or martensite or the ferrite + bainite two-phase structure is formed, so that the HIC resistance is deteriorated. . On the other hand, when the cooling stop temperature exceeds 600 ° C., the ferrite transformation at the time of reheating is not completed, pearlite is precipitated and the HIC resistance is deteriorated, and the effect of transformation strengthening by bainite transformation is not sufficient, and the strength is high. descend. The cooling stop temperature is preferably 400 to 600 ° C. from the viewpoint of increasing the driving force of the ferrite transformation at the time of reheating and sufficiently obtaining the effect of precipitation strengthening by precipitates at the time of ferrite transformation.

Immediately after the accelerated cooling described above, reheating is performed at a temperature rising rate of 0.5 ° C./sec or more to a temperature equal to or higher than the cooling stop temperature and to a temperature of 550 to 700 ° C. This process is an important manufacturing condition in the present invention. Fine precipitates that contribute to strengthening of the ferrite phase are deposited simultaneously with the ferrite transformation during reheating. In order to simultaneously strengthen the ferrite phase by the fine precipitates and soften the bainite phase, and obtain a structure having a small strength difference between the ferrite phase and the bainite phase, immediately after the accelerated cooling, the cooling stop temperature is not less than 550 to 700. It is necessary to reheat to the temperature range of ° C. In this reheating, it is desirable to raise the temperature to 50 ° C. or higher than the cooling stop temperature.
If the heating rate is less than 0.5 ° C./sec, it takes a long time to reach the target reheating temperature, so that the dispersion of fine precipitates cannot be obtained and sufficient strength cannot be obtained. , Manufacturing efficiency deteriorates. Moreover, in order to suppress the deterioration of toughness, it is effective to finely and uniformly disperse and precipitate by suppressing the coarsening of the carbide during the temperature rise. From this viewpoint, the temperature rise rate is 3 ° C./sec or more. It is preferable that

  The reheating temperature is equal to or higher than the cooling stop temperature to double tempering. Further, if the reheating temperature is less than 550 ° C., sufficient precipitation strengthening by fine precipitates cannot be achieved, and ferrite transformation is not completed, and untransformed austenite is transformed into pearlite during subsequent cooling, resulting in deterioration of HIC resistance. On the other hand, if the reheating temperature exceeds 700 ° C., the precipitate becomes coarse and sufficient strength cannot be obtained. There is no need to set the temperature holding time at the reheating temperature. Therefore, it may be cooled immediately after reaching the reheating temperature. The cooling rate is appropriately selected so that fine precipitates are continuously deposited, and air cooling is particularly desirable. When the temperature is maintained at the reheating temperature, if the temperature is maintained for more than 30 minutes, the precipitates are coarsened and the strength may be reduced.

  FIG. 2 shows an example of equipment suitable for manufacturing the high-strength steel sheet of the present invention. In the rolling line 1, a hot rolling mill 3, an accelerated cooling device 4, a hot leveler 5, and an inline induction heating device 6 for reheating the steel plate after accelerated cooling are arranged from the upstream side toward the downstream side. Since the inline induction heating device 6 is installed on the same line as the hot rolling mill 3 and the accelerated cooling device 4, the steel plate 2 after the completion of rolling and cooling can be quickly reheated. That is, the steel plate 2 after rolling and accelerated cooling can be immediately reheated to the cooling stop temperature or higher and to 550 to 700 ° C. without excessive cooling from the cooling stop temperature.

The induction heating apparatus used in FIG. 2 is particularly preferable because temperature control is easier than in a soaking furnace, the equipment cost is relatively low, and the cooled steel sheet can be heated quickly.
Also, by arranging a plurality of induction heating devices in series, even if the line speed and the type / size of the steel sheet are different, the temperature increase rate can be set simply by arbitrarily setting the number of induction heating devices to be energized. The reheating temperature can be freely controlled.
Since the cooling rate after reheating is arbitrary, it is not necessary to install special equipment downstream of the heating device. As a heating device, a gas combustion furnace capable of rapid heating of a steel plate may be used instead of the in-line induction heating device 6.

  The high-strength steel sheet of the present invention is formed into a tubular shape by press bend forming, roll forming, UOE forming, etc., and then welded (further expansion or the like is performed if necessary), which is suitable for transportation of crude oil and natural gas. High-strength steel pipes for line pipes (UOE steel pipes, ERW steel pipes, spiral steel pipes, etc.) excellent in HIC resistance and weld heat affected zone toughness can be produced. For example, in UOE steel pipe, the edge part of a steel plate is grooved, and after forming into a ring shape by C press, U press, O press, the groove part is welded by temporary welding and inner / outer surface welding, and undergoes a pipe expanding process. Manufactured.

Steels having the chemical components shown in Table 1 (steel types A to Y) were made into slabs by a continuous casting method, and No. 2 shown in Tables 2 and 3 were used. 1-No. 33 thick steel plates were produced.
After heating the slab to a predetermined thickness by hot rolling, it was immediately cooled using a water-cooled accelerated cooling facility, and then reheated using an induction heating furnace or a gas combustion furnace. The induction furnace was installed on the same line as the accelerated cooling equipment.
The metal structure of the obtained steel sheet was observed with an optical microscope and a transmission electron microscope (TEM). For the metal structure, the central portion of the plate thickness and the t / 4 position were observed with an optical microscope, the area fraction of the ferrite phase and the bainite phase was measured from the photographed image by image processing, and the area fraction of each phase in five fields of view. Was the volume fraction. Moreover, the component of the deposit was analyzed by energy dispersive X-ray spectroscopy (EDX). In addition, the tensile properties, HIC resistance, and weld heat affected zone (HAZ) toughness of each steel plate were measured. The results are shown in Tables 2 and 3 together with the production conditions.

Tensile properties were measured by performing a tensile test using a full thickness test piece in the vertical direction of rolling as a tensile test piece, and measuring the tensile strength. With respect to the weld heat affected zone (HAZ) toughness, a Charpy test was performed using a test piece to which a heat history corresponding to a maximum heating temperature of 1400 ° C. and a heat input of 40 kJ / cm was added by a reproducible heat cycle apparatus. HIC resistance is determined by conducting an HIC test with an immersion time of 96 hours in accordance with NACE Standard TM-02-84. If no cracks are observed, it is determined that the HIC resistance is good. Evaluated as “×”.
In the performance evaluation of this example, considering the manufacturing variation, the tensile strength is 580 MPa or more (APIX 70 grade or more), the HAZ toughness is ductile-brittle transition temperature (vTrs) of 0 ° C. or less, the HIC resistance is not cracked, Each was accepted.

In Table 2 and Table 3, no. Examples 1 to 18 are examples of the present invention, all having good HIC resistance, a tensile strength of 580 MPa or more, and a ductile-brittle transition temperature (vTrs) of a weld heat affected zone of 0 ° C. or less. Further, it was observed that fine composite carbides having a particle size of less than 10 nm containing one or two of Nb, V, and Ti were dispersed and precipitated at a density of 2 × 10 ³ particles / μm ² or more. It was.
On the other hand, no. Nos. 19 to 23 are comparative examples in which the chemical components satisfy the conditions of the present invention, but the production method does not satisfy the conditions of the present invention, and all of them are insufficiently dispersed and precipitated of fine carbides, and sufficient tensile strength is obtained. Not. No. No. 19 has a low slab heating temperature and insufficient solid solution of carbides required for fine dispersion precipitation. No. 20 and no. In No. 21, since accelerated cooling does not satisfy the conditions of the present invention, a two-phase structure of ferrite phase + bainite phase cannot be obtained, and dispersion and precipitation of fine carbides is insufficient. Furthermore, since island-like martensite (MA) and pearlite are precipitated, the HIC resistance is inferior. No. No. 22 has a slow reheating temperature rise rate, so that the dispersion and precipitation of fine carbides is insufficient, and sufficient tensile strength is not obtained. No. Since No. 23 has a low reheating temperature, fine carbide dispersion and precipitation are insufficient, and sufficient tensile strength is not obtained. Moreover, since pearlite precipitates, the HIC resistance is inferior.
No. Nos. 24-33 are inferior in either HIC resistance or HAZ toughness because the chemical components do not satisfy the conditions of the present invention.

DESCRIPTION OF SYMBOLS 1 Rolling line 2 Steel plate 3 Hot rolling mill 4 Accelerated cooling device 5 Hot leveler 6 In-line type induction heating device

Claims (4)

In mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 1.8%, P: 0.01% or less, S: 0.002 %: Ca: 0.0005 to 0.005%, Nb: 0.05 to 0.15%, Al: 0.01 to 0.08%, and V: 0.005 to 0.15 %, Ti: One or two of 0.005 to 0.04%, the balance being Fe and inevitable impurities, and the ratio of the amount of C in atomic% and the total amount of Nb, V and Ti C / (Nb + V + Ti ) is 1.0 to 5.0 is the following (1) CP value represented by the formula (wt%) 0.98 or less, _{P CM} value represented by the following formula (2) ( (Mass%) has a component composition of 0.15 or less,
The metal structure is a substantial two-phase structure in which the sum of the ferrite phase and the bainite phase is 95% or more by volume fraction, and carbides containing one or two of Nb, V, and Ti are dispersed and precipitated. A high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness, characterized by having a tensile strength of 580 MPa or more.
CP = 4.46C (%) + 2.37Mn (%) / 6+ {1.74Cu (%) + 1.7Ni (%)} / 15+ {1.18Cr (%) + 1.95Mo (%) + 1.74V (% )} / 5 + 22.36P (%) (1)
P _CM = C (%) + Si (%) / 30 + Mn (%) / 20 + Cu (%) / 20 + Ni (%) / 60 + Cr (%) / 20 + Mo (%) / 15 + V (%) / 10 + B (%) * 5 (5) 2)
However, in the formulas (1) and (2), the element not added is 0.   Further, by mass%, Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less, Mo: 0.50% or less, B: 0.005% or less The high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness according to claim 1, comprising at least a seed. It is a manufacturing method of the high strength steel plate for line pipes according to claim 1 or 2,
A steel having the component composition according to claim 1 or 2 is heated to a temperature of 1000 to 1300 ° C and hot-rolled at a rolling end temperature of Ar ₃ temperature or higher, and then a cooling rate of 5 ° C / sec or higher. Accelerating cooling to 300 to 600 ° C., and immediately after that, at a temperature rising rate of 0.5 ° C./sec or more, reheating to 550 to 700 ° C. is performed at or above the cooling stop temperature. A method for producing high-strength steel sheets for line pipes with excellent HIC characteristics and weld heat-affected zone toughness.   A high-strength steel pipe for a line pipe excellent in HIC resistance and weld heat affected zone toughness, characterized by being manufactured using the steel plate according to claim 1 or 2.

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