JP4419744B2 - 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 having an API standard X65 grade or higher suitable for the manufacture of steel pipes and the like, and particularly high strength for line pipes excellent in hydrogen-induced crack resistance (HIC resistance) and weld heat affected zone toughness. It is related with a steel plate and its manufacturing method.

  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, around non-metallic inclusions such as MnS in the steel and hard second phase structure. It diffuses and accumulates on the surface and cracks are caused by its internal pressure. Furthermore, along with the development of cold regions, even when steel sheets are used at temperatures below -30 ° C, there is a tendency to demand strict performance above a certain level. There are cases where a certain level or more is required.

  In order to prevent such hydrogen-induced cracking, by adding an appropriate amount of Ca or Ce with respect to the amount of S, the formation of acicular MnS is suppressed, and the form is formed into finely dispersed spherical inclusions with a small stress concentration. A manufacturing method of steel for line pipes that is excellent in HIC resistance and that suppresses the generation and propagation of cracks by changing is known (see, for example, Patent Document 1). In addition, islands that are the starting point of cracks in the central segregation part due to reduction of elements with high segregation tendency (C, Mn, P, etc.), soaking in the slab heating stage, and accelerated cooling during transformation during cooling Steels excellent in HIC resistance that suppress the formation of martensite and hardened structures such as martensite and bainite, which are propagation paths of cracks, are known (see, for example, Patent Document 2 and Patent Document 3). In addition, regarding X80 grade high-strength steel sheets with excellent HIC resistance, while controlling the form of inclusions by adding Ca at low S, suppressing central segregation as low C and low Mn, the accompanying strength reduction is Cr, A method of supplementing by adding Mn, Ni or the like and accelerated cooling is known (see, for example, Patent Document 4, Patent Document 5, and Patent Document 6).

  However, all the methods for improving the above-mentioned HIC resistance are for the center segregation part. High-strength steel sheets exceeding the X65 grade such as API X80 grade are often manufactured by accelerated cooling or direct quenching, so the steel plate surface portion with a high cooling rate is hardened compared to the inside, and hydrogen-induced cracking occurs from the vicinity of the surface. . Moreover, the microstructure of these high-strength steel sheets obtained by accelerated cooling is a relatively high cracking susceptibility of bainite or acicular ferrite not only to the surface but also to the inside. Even in this case, it is difficult to eliminate HIC starting from sulfide or oxide inclusions in high strength steel of about API X80 grade. Therefore, when the HIC resistance of these high-strength steel plates is a problem, it is necessary to take measures against HIC on the surface portion of the steel plate or HIC starting from sulfide or oxide inclusions.

On the other hand, as a high strength steel excellent in HIC resistance that does not contain block-like bainite or martensite with high cracking susceptibility, high strength with excellent HIC resistance of API X80 grade, which is a ferrite-bainite two-phase structure Steel materials are known (see, for example, Patent Document 7). In addition, high strength steel is known that improves the SCC (SSCC) resistance and HIC resistance by making the microstructure a ferrite single phase structure, and uses carbide precipitation strengthening obtained by adding a large amount of Mo or Ti. (For example, see Patent Document 8 and Patent Document 9).
Japanese Patent Laid-Open No. 54-110119 JP 61-60866 A JP-A-61-165207 JP-A-5-9575 JP-A-5-271766 JP 7-173536 A JP 7-216500 A Japanese Patent Laid-Open No. 61-227129 JP-A-7-70697

  However, the bainite structure of high-strength steel described in Patent Document 7 and the like is a structure that is not as high as block bainite and martensite but is relatively high in cracking sensitivity. Since it is necessary to improve the HIC resistance, 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 7 and the like. On the other hand, the ferrite phase described in Patent Document 8, Patent Document 9 and the like has a highly ductile structure and extremely low susceptibility to cracking, so that the HIC resistance is greatly improved compared to steels having a bainite structure or an acicular ferrite structure. The However, since the strength of the ferrite single phase is low, the steel described in Patent Document 8 is strengthened by using a steel containing a large amount of C and Mo to precipitate a large amount of carbides. Then, Ti-added steel is wound around a steel strip at a specific temperature, and the strength is increased by utilizing precipitation strengthening of TiC. However, in order to obtain a ferrite structure in which Mo carbides described in Patent Document 8 are dispersed, it is necessary to perform cold working after quenching and tempering, and further tempering again, which not only increases the manufacturing cost but also increases Mo. Since the particle size of the carbide is as large as about 0.1 microns 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 the carbide. In addition, TiC used in the high-strength steel described in Patent Document 9 is finer than Mo carbide and is an effective carbide for precipitation strengthening, but it is likely to be coarsened under the influence of temperature during precipitation. Regardless, no countermeasures against precipitate coarsening have been taken. Therefore, 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.

  Therefore, the object of the present invention is to solve such problems of the prior art, and is a high-strength steel sheet for line pipes of API X65 grade or higher, in which HIC in the central segregation part and HIC generated from the vicinity of the surface and inclusions Thus, it is to provide a high-strength steel sheet for line pipes having excellent HIC resistance and excellent weld heat-affected zone toughness without adding a large amount of alloy elements at a low cost.

The features of the present invention for solving such problems are as follows.
(1), in mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 2%, P: 0.01% or less, S: 0 0.002% or less, Mo: 0.05 to 0.5%, Ti: 0.005 to 0.025%, Al: 0.07% or less, N: 0.004 to 0.006%, the balance Is composed of Fe and inevitable impurities , and C / (Mo + Ti), which is the ratio of the amount of C in atomic% to the total amount of Mo and Ti, is 0.5 to 3, and the amount of Ti and the amount of N in mass% The Ti / N ratio is 2 to 4, the metal structure has a total volume fraction of ferrite and bainite of 95% or more , and precipitates containing Ti and Mo are dispersed and precipitated. A high-strength steel sheet for line pipes with excellent HIC resistance and weld heat-affected zone toughness.
(2) Further, by mass%, Nb: 0.005 to 0.05% and / or V: 0.005 to 0.1%, C amount in atomic% and Mo, Ti, Nb, C / (Mo + Ti + Nb + V), which is a ratio of the total amount of V, is 0.5 to 3, and a composite precipitate containing Ti, Mo, Nb and / or V is dispersedly precipitated. A high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness according to (1).
(3) By mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 2%, P: 0.01% or less, S: 0 0.002% or less, Ti: 0.005 to 0.025%, Al: 0.07% or less, N: 0.004 to 0.006%, Nb: 0.005 to 0.05% and / or Or V: 0.005 to 0.10%, the balance is Fe and inevitable impurities , and C / (Ti + Nb + V), which is the ratio of the amount of C in atomic% and the total amount of Ti, Nb, and V, 0.5-3, Ti / N ratio of Ti amount and N amount in mass% is 2-4, and the total volume fraction of ferrite and bainite is 95% or more in metal structure A composite carbide containing Ti and Nb and / or V is dispersed and precipitated, and the HIC resistance and welding heat High strength steel plate for line pipe superior in sound zone toughness.
(4) Further, by mass%, Cu: 0.5% or less, Ni: 0.5% or less, Cr: 0.5% or less, Ca: 0.0005 to 0.003%, Mg: 0.005 % For line pipes having excellent HIC resistance and weld heat affected zone toughness as set forth in any one of (1) to (3) Strength steel plate.
(5) After heating the steel having the component composition according to any one of (1) to (4) to a temperature of 1000 to 1300 ° C. and hot rolling at a rolling finish temperature not lower than the Ar 3 temperature, 5 ° C. HIC resistance characteristics, characterized in that accelerated cooling to 300 to 600 ° C. at a cooling rate of at least / s, and immediately reheating to 550 to 700 ° C. at a temperature rising rate of at least 0.5 ° C./s A method for producing high-strength steel sheets for line pipes with excellent weld heat-affected zone toughness.
(6) A high-strength steel pipe excellent in HIC resistance and weld heat-affected zone toughness, characterized by being manufactured using the steel plate according to any one of (1) to (4).

  According to the present invention, it is possible to manufacture a steel sheet having high strength equal to or higher than API X65 grade and having excellent HIC resistance and weld heat affected zone toughness without adding a large amount of alloy elements. . For this reason, steel pipes, such as an electric resistance welded steel pipe, a spiral steel pipe, and a UOE steel pipe, having excellent characteristics can be manufactured.

  The present inventors diligently studied the microstructure of the steel material and the manufacturing method of the steel sheet in order to achieve both improved HIC resistance and high strength, and further improve the weld heat affected zone toughness. As a result, to achieve both high strength and HIC resistance, it is most effective to make the microstructure a ferrite + bainite two-phase structure with a small difference in strength between the ferrite structure and the bainite structure. By performing a manufacturing process of accelerated cooling and subsequent reheating, strengthening of the ferrite phase, which is a soft phase due to fine precipitates containing Ti, Mo, etc., and fine precipitates containing Ti, Nb, V, etc., and a hard phase As a result, the bainite phase was softened and a ferrite + bainite two-phase structure with a small strength difference could be obtained. And the knowledge that precipitation strengthening by a carbide | carbonized_material can be utilized to the maximum was acquired by optimizing the addition amount of Mo and Ti with respect to C, and the addition amount of Ti, Nb, and V with respect to C. Also, when the steel material contains Mo, if Nb and / or V are added in combination, the strength of the ferrite phase is increased by dispersing and precipitating Ti, Mo, and Nb and / or V precipitates. It was found that the precipitation strengthening by carbides can be utilized to the maximum by optimizing the addition amount of Mo, Ti, Nb and V to C.

  The present invention is excellent in HIC resistance having a two-phase structure of a precipitate containing Ti, Mo and the like, a ferrite phase in which a precipitate containing Ti, Nb, V and the like is dispersed and a bainite phase. The high strength steel sheet for line pipes and the manufacturing method thereof, and the steel sheet manufactured in this way has increased hardness in the surface layer portion such as a steel sheet of bainite or acicular ferrite structure obtained by conventional accelerated cooling or the like. Therefore, HIC from the surface layer does not occur. Furthermore, since the two-phase structure of the ferrite phase and the bainite phase having a small strength difference has an extremely high resistance to cracking, it is possible to suppress HIC from the central part of the steel sheet and inclusions. Furthermore, in the weld heat affected zone that reaches 1350 ° C. or higher during welding, the prior austenite grains in the weld heat affected zone are refined by finely dispersing TiN at high temperatures, and the toughness of the weld heat affected zone is improved. Obtained knowledge.

  Hereinafter, the high-strength steel sheet for line pipes of the present invention will be described in detail. First, the structure of the high-strength steel sheet for line pipes of the present invention will be described.

The metal structure of the steel sheet of the present invention is substantially a ferrite + bainite two-phase structure.
Since the ferrite phase is rich in ductility and has low cracking susceptibility, high HIC resistance can be realized. The bainite phase has excellent strength toughness. 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. However, in the present invention, both the HIC resistance and high strength can be achieved by adjusting the strength of the ferrite phase and the bainite phase to reduce the strength difference between the two. When two or more different metal structures such as martensite and pearlite are mixed in the ferrite + bainite two-phase structure, HIC is likely to occur due to hydrogen accumulation and stress concentration at the heterophase interface. The smaller the fraction of the structure other than the phase, the better. However, when the volume fraction of the structure other than the ferrite phase and the bainite phase is low, the influence can be ignored. Therefore, other metal structures of less than 5% in total volume fraction, that is, one kind of martensite, pearlite, or the like You may contain 2 or more types. The bainite fraction is not particularly defined, but is preferably 10% or more from the viewpoint of securing the toughness of the base material and 80% or less from the viewpoint of HIC resistance. More preferably, it is 20 to 60%.

  Next, regarding the precipitates that are dispersed and precipitated in the ferrite phase in the present invention, the case where the steel sheet contains Mo will be described as a first embodiment, and the case where no steel is contained will be described as a second embodiment.

  In the steel sheet of the present invention in the first embodiment, the ferrite phase is strengthened by the precipitation containing Mo and Ti as a basis in the ferrite phase, and the strength difference between ferrite and bainite is reduced. Excellent anti-HIC characteristics can be obtained. Since this precipitate is extremely fine, it has no influence on the HIC resistance. Mo and Ti are elements that form carbides in steel, and it has been conventionally practiced to strengthen steel by precipitation of MoC and TiC. However, in the present invention, Mo and Ti are added together to form Mo and Ti. It is a feature that a greater strength improvement effect can be obtained by finely precipitating the composite carbide containing the above in steel as compared with the case of precipitation strengthening of MoC and / or TiC. This unprecedented strength improvement effect is due to the fact that the composite carbide containing Mo and Ti as a basis is stable and has a slow growth rate, so that an extremely fine precipitate having a particle size of less than 10 nm can be obtained. .

  When the composite carbide containing Mo and Ti as a base is composed of only Mo, Ti, and C, the total amount of Mo and Ti and the amount of C are combined in an atomic ratio of about 1: 1. It is very effective in increasing strength. Further, in the present invention, it has been found that by adding Nb and / or V in combination, the precipitate becomes a composite carbide containing Mo, Ti, Nb and / or V, and the same precipitation strengthening can be obtained. It was. Moreover, although this fine carbide mainly precipitates in the ferrite phase, it may also precipitate from the bainite phase depending on chemical components and production conditions. In this case, further strengthening is possible, but if the hardness difference between the ferrite phase and the bainite phase is HV70 or less, the HIC resistance is not affected.

In order to obtain a high-strength steel sheet having a yield strength of 448 MPa or more (API × 65 grade or more), it is preferable to deposit 2 × 10 ³ pieces / μm ³ or more. The form of precipitation may be random or in line and is not particularly defined. Moreover, when it contains precipitates other than the composite carbide mainly composed of Mo and Ti, the effect of increasing the strength by the composite carbide of Mo and Ti is not impaired and the HIC resistance is not deteriorated, but less than 10 nm. The number of precipitates is preferably 95% or more of the total number of precipitates excluding TiN. The average particle size of the fine composite carbide precipitates is obtained by subjecting a photograph taken with a transmission electron microscope (TEM) to image processing, and obtaining the diameter of a circle having the same area as each precipitate for each composite carbide. , And can be obtained as an average value of their diameters.

  Further, in the steel plate of the present invention in the second embodiment, the ferrite phase is strengthened by the composite carbide containing Ti and Nb and / or V being dispersed and precipitated in the ferrite phase, and the strength between ferrite and bainite. Since the difference is reduced, excellent HIC resistance can be obtained. Since this precipitate is extremely fine, it has no influence on the HIC resistance. Ti, Nb, and V are elements that form carbides in steel, and strengthening of steel by precipitation of these carbides has been conventionally performed, but conventionally, by cooling process and isothermal holding after hot rolling. There is a method of using precipitation from supersaturated ferrite during ferrite transformation from austenite, or quenching after rolling to make the structure martensite or bainite, and then precipitating carbide in martensite or bainite by tempering. It was used. On the other hand, in the present invention, the carbide is precipitated by utilizing the ferrite transformation in the reheating process from the bainite transformation region. According to this method, the ferrite transformation proceeds very fast, and fine composite carbide containing Ti and Nb and / or V with a grain size of less than 20 nm is precipitated at the transformation interface. It is characterized in that a greater strength improvement effect can be obtained.

  The composite carbide containing Ti and Nb and / or V is a compound in which the total amount of Ti, Nb, and V and the amount of C are combined in an atomic ratio of about 1: 1, which increases the strength. It is very effective and can be obtained by manufacturing a steel sheet using the manufacturing method of the present invention to the steel having the components described below. Although this fine carbide is mainly precipitated in the ferrite phase, it may be precipitated from the bainite phase depending on chemical components and production conditions. In this case, further strengthening is possible, but if the hardness difference between the ferrite phase and the bainite phase is HV70 or less, the HIC resistance is not affected. Note that the composite carbide containing Ti and Nb and / or V is dispersed and precipitated simultaneously with the composite carbide containing Mo and Ti as a basis even in the first embodiment in which the steel contains Mo. In some cases.

The number of precipitates of less than 20 nm is preferably deposited at 2 × 10 ³ pieces / μm ³ or more in order to obtain a high-strength steel sheet having a yield strength of 448 MPa or more (API × 65 grade or more). The form of precipitation may be random or in a row and is not particularly defined. When the high-strength steel sheet for line pipes of the present invention contains precipitates other than composite carbides containing Ti and Nb and / or V, the effect of increasing strength by the composite carbides is not impaired, and HIC resistance However, the number of precipitates of less than 20 nm is preferably 95% or more of the total number of precipitates excluding TiN. The average particle size of the fine composite carbide precipitates is obtained by subjecting a photograph taken with a transmission electron microscope (TEM) to image processing, and obtaining the diameter of a circle having the same area as each precipitate for each composite carbide. , And can be obtained as an average value of their diameters.

  In the present invention, composite carbides mainly composed of Mo and Ti, and composite carbides including Ti, Nb, V, and the like, which are precipitates dispersed and precipitated in the steel sheet, are produced by the production method of the present invention on steel having the components described below. It can obtain by manufacturing a steel plate using.

  Next, chemical components of the high-strength steel sheet for line pipe used in the present invention will be described. Unless otherwise specified in the following description, all units shown in% are% by mass.

  C: Set to 0.02 to 0.08%. C is an element that contributes to precipitation strengthening as a carbide. However, if it is less than 0.02%, sufficient strength cannot be secured, and if it exceeds 0.08%, toughness and HIC resistance are deteriorated. 0.02 to 0.08%. More preferably, it is 0.03 to 0.06%.

  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, and if it exceeds 0.5%, the toughness and weldability are deteriorated, so the Si content is 0.01 to 0.00. Specify 5%. More preferably, it is 0.01 to 0.3%.

  Mn: 0.5 to 2%. Mn is added for strength and toughness, but if it is less than 0.5%, the effect is not sufficient, and if it exceeds 2%, the weldability and HIC resistance deteriorate, so the Mn content is 0.5-2%. Stipulate. Preferably, it is 0.5% or more and less than 1.5%.

  P: 0.01% or less. Since P is an inevitable impurity element that deteriorates weldability and HIC resistance, the upper limit of the P content is specified to be 0.01%.

  S: Set to 0.002% or less. In general, S is preferably as small as possible because it becomes MnS inclusions in steel and deteriorates the HIC resistance. However, since there is no problem if it is 0.002% or less, the upper limit of the S content is defined as 0.002%.

  Ti: 0.005 to 0.025%. Ti is an important element in the present invention. By adding 0.005% or more, a composite precipitate is formed together with Mo, Nb and / or V, which greatly contributes to an increase in strength. However, if added over 0.025%, the weld heat-affected zone toughness is deteriorated, so the Ti content is specified to be 0.005 to 0.025%. Furthermore, when it is less than 0.02%, more excellent toughness is exhibited. For this reason, when adding Nb and / or V, it is preferable to make Ti content into 0.005 to less than 0.02%.

  Al: 0.07% or less. Al is added as a deoxidizer, but if it exceeds 0.07%, the cleanliness of the steel is lowered and the HIC resistance is deteriorated, so the Al content is specified to be 0.07% or less. Preferably, it is 0.01 to 0.07%.

  N: Set to 0.004 to 0.006%. N forms precipitates with Ti and finely disperses in the high temperature region of the weld heat affected zone reaching 1350 ° C or higher, thereby refining the old austenite grains in the weld heat affected zone and improving the toughness of the weld heat affected zone. A big contribution. If less than 0.004%, the effect is not sufficient, and if added over 0.006%, the precipitates become coarse, leading to toughness deterioration of the weld heat affected zone and also causing slab cracking in the steelmaking stage. The amount is 0.004 to 0.006%.

  Ti / N: 2-4. By defining the Ti / N amount, which is the ratio of the Ti amount and the N amount in mass%, within an appropriate range, TiN is finely dispersed, and refinement of the prior austenite grains in the weld heat affected zone is achieved. The toughness of the affected area is improved. When the amount of Ti / N is less than 2, the effect is not sufficient, and when it exceeds 4, the precipitate becomes coarse and the toughness of the weld heat affected zone deteriorates, so the value of Ti / N is defined as 2 to 4.

  The steel plate of the present invention in the first embodiment contains Mo.

  Mo: 0.05 to 0.5%. Mo is an important element in the case of precipitating a composite carbide containing Mo and Ti as a base, and when added, 0.05% or more is included, so that pearlite transformation at the time of cooling after hot rolling is performed. While suppressing, a fine composite precipitate with Ti is formed, which greatly contributes to an increase in strength. However, if added over 0.5%, a hardened phase such as martensite is formed and the HIC resistance is deteriorated, so the Mo content is specified to be 0.05 to 0.5%. Preferably, it is 0.05% to 0.3%.

  C / (Mo + Ti), which is the ratio of the amount of C and the total amount of Mo and Ti, is 0.5-3. Strengthening in the case of containing Mo, which is the first embodiment of the present invention, is due to precipitates (mainly carbides) containing Ti and Mo. In order to effectively use the precipitation strengthening by this composite precipitate, the relationship between the amount of C and the amounts of Mo and Ti which are carbide forming elements is important, and these elements should be added in an appropriate balance. Thus, a thermally stable and very fine composite precipitate can be obtained. At this time, when the value of C / (Mo + Ti) represented by the content of atomic% of each element is less than 0.5 or more than 3, the amount of any element is excessive, and the HIC resistance due to the formation of a hardened structure In order to cause deterioration of characteristics and toughness, the value of C / (Mo + Ti) is regulated to 0.5-3. However, each element symbol is the content of each element in atomic%. In addition, when content of the mass% is used, it is represented by (C / 12.01) / (Mo / 95.9 + Ti / 47.9). The value of C / (Mo + Ti) is specified to be 0.5-3. When the value of C / (Mo + Ti) is 0.7-2, a finer precipitate having a particle size of 5 nm or less is obtained, which is more preferable.

  In the case of the first embodiment containing Mo, Nb and / or V shown below may be contained for the purpose of further improving the strength and weld zone toughness of the steel sheet. In the case of the second embodiment of the present invention that does not contain Mo, the steel sheet of the present invention contains Nb and / or V in order to form a composite precipitate together with Ti. Even when Nb and / or V is contained, the metal structure is substantially a two-phase structure of ferrite and bainite.

  Nb: 0.005 to 0.05%. Nb improves toughness by refining the structure, but forms a composite precipitate with Ti and Mo, Ti and V, and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.05%, the toughness of the weld heat affected zone deteriorates, so the Nb content is specified to be 0.005 to 0.05%.

  V: Set to 0.005 to 0.1%. V, like Nb, forms a composite precipitate with Ti and Mo, Ti and Nb, and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.1%, the toughness of the weld heat affected zone deteriorates, so the V content is specified to be 0.005 to 0.1%.

  In the first embodiment of the present invention containing Mo, when Nb and / or V is contained, the ratio of the amount of C and the total amount of Mo, Ti, Nb, and V is C / (Mo + Ti + Nb + V): Is 0.5-3. Strengthening in the case of containing Mo depends on precipitates containing Ti and Mo, but in the case of containing Nb and / or V, composite precipitates (mainly carbides) containing them are formed. At this time, when the value of C / (Mo + Ti + Nb + V) represented by the content of atomic% of each element is less than 0.5 or more than 3, the amount of any element is excessive, and HIC resistance due to formation of a hardened structure In order to cause deterioration of characteristics and toughness, the value of C / (Mo + Ti + Nb + V) is regulated to 0.5-3. However, each element symbol is a content in atomic%. In addition, when content of the mass% is used, the value of (C / 12.01) / (Mo / 95.9 + Ti / 47.9 + Nb / 92.91 + V / 50.94) is specified to 0.5-3. . More preferably, it is 0.7-2, and a finer precipitate having a particle size of 5 nm or less is obtained.

  Further, in the case of the second embodiment in which the high-strength steel sheet of the present invention does not contain Mo, C / (Ti + Nb + V), which is the ratio of the amount of C in atomic% and the total amount of Ti, Nb, and V, is 0. .5 to 3. The increase in strength in this case is due to the precipitation of fine carbides containing Ti and Nb and / or V. In order to effectively use the precipitation strengthening by this fine carbide, the relationship between the amount of C and the amounts of Ti, Nb, and V which are carbide forming elements is important, and these elements are added in an appropriate balance. By doing so, a thermally stable and very fine composite carbide can be obtained. At this time, when the value of C / (Ti + Nb + V) represented by the atomic% content of each element is less than 0.5 or more than 3, the amount of any element is excessive, and HIC resistance due to formation of a hardened structure In order to cause deterioration of characteristics and toughness, the value of C / (Ti + Nb + V) is regulated to 0.5-3. However, each element symbol is the content of each element in atomic%. In addition, when content of the mass% is used, the value of (C / 12.01) / (Ti / 47.9 + Nb / 92.91 + V / 50.94) is specified to 0.5-3. When the value of C / (Ti + Nb + V) is 0.7-2, it is more preferable because a finer precipitate can be obtained.

  In the present invention, for the purpose of further improving the strength and HIC resistance of the steel sheet, one or more of Cu, Ni, Cr, Ca and Mg shown below may be contained.

  Cu: 0.5% or less. Cu is an element effective for improving toughness and increasing strength. In order to acquire the effect, it is preferable to add 0.1% or more. However, if a large amount is added, weldability deteriorates.

  Ni: 0.5% or less. Ni is an element effective for improving toughness and increasing strength. In order to obtain the effect, it is preferable to add 0.1% or more. However, if added in a large amount, the HIC resistance is lowered, so when added, the upper limit is 0.5%.

  Cr: 0.5% or less. Cr, like Mn, is an element effective for obtaining sufficient strength even at low C. In order to acquire the effect, it is preferable to add 0.1% or more. However, if a large amount is added, the weldability is deteriorated.

  Ca: 0.0005 to 0.003%. Ca is an element effective for improving the HIC resistance by controlling the form of sulfide inclusions, but the effect is not sufficient if it is less than 0.0005%, and the effect is saturated even if added over 0.003%. However, since the HIC resistance is deteriorated due to a decrease in the cleanliness of the steel, the Ca content is specified to be 0.0005 to 0.003% when added.

  Mg: 0.005% or less. Mg is an element that contributes to improvement of the weld heat-affected zone toughness. In order to obtain the effect, it is preferable to add 0.0005% or more, but if added over 0.005%, the weld heat affected zone toughness is deteriorated conversely. And

  Moreover, it is preferable to prescribe | regulate the upper limit of Ceq defined by the following (a) formula from a viewpoint of weldability according to a strength level. When the yield strength is 448 MPa or more, Ceq is 0.28 or less, when the yield strength is 482 MPa or more, Ceq is 0.32 or less, and when the yield strength is 551 MPa or more, Ceq is 0.36 or less. By making it, good weldability can be secured.

Ceq = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (a)
However, the element symbol of the formula (a) indicates mass% of each contained element.

  In the case of the second embodiment in which the steel sheet of the present invention does not contain Mo, the following formula (b) may be used instead of formula (a).

Ceq = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + V) / 5 (b)
However, the element symbol of the formula (b) indicates mass% of each contained element.

  In addition, about the steel material of this invention, there is no board thickness dependence of Ceq in the range of about 10 mm to 30 mm, and it can design with the same Ceq to about 30 mm.

  The remainder other than the above consists essentially of Fe, and it is possible to add a trace amount of elements that do not impair the effects of the present invention, including inevitable impurities. For example, you may add REM, W, and Zr 0.02% or less, respectively.

  Next, the manufacturing method of the high strength steel plate for line pipes of this invention is demonstrated.

  FIG. 1 is a diagram schematically showing the tissue control method of the present invention. By performing accelerated cooling (C) from the austenite region (A) of Ar 3 or higher to the bainite transformation region (B), a mixed structure of untransformed austenite 11 and bainite 12 is obtained from the austenite single phase 10. By reheating (D) to the ferrite region (E) immediately after cooling, the austenite 11 is transformed into ferrite, and fine precipitates are dispersed and precipitated in the ferrite phase. 13 On the other hand, the bainite phase 12 is tempered to become tempered bainite 14. The tempered bainite 14 may be strengthened by precipitation by dispersing fine precipitates. Hereinafter, the tissue control method will be specifically described in detail.

  The high-strength steel sheet for line pipes of the present invention uses steel having the above-described composition, and is hot-rolled at a heating temperature of 1000 to 1300 ° C. and a rolling end temperature of Ar 3 temperature or higher, and then cooled at 5 ° C./s or higher. It is cooled to 300 to 600 ° C. at a speed, and immediately after cooling, it is reheated to a temperature of 550 to 700 ° C. at a temperature rising rate of 0.5 ° C./s or more, so that the metal structure becomes a two phase structure of ferrite and bainite. It can be produced as a composite structure in which fine composite carbides mainly composed of Mo and Ti and fine composite carbides containing Ti and Nb and / or V are dispersed and precipitated to soften the bainite phase. Here, the temperature is the average temperature of the steel sheet. The average temperature is obtained by calculation based on the surface temperature of the slab or steel plate, taking into account parameters such as plate thickness and thermal conductivity. Moreover, a cooling rate is an average cooling rate which divided the temperature difference required for cooling to the cooling end temperature (300-600 degreeC) after the completion | finish of hot rolling by the time required to perform the cooling. The temperature increase rate is an average temperature increase rate obtained by dividing the temperature difference required for reheating up to the reheating temperature (550 to 700 ° C.) by the time required for reheating after cooling. Hereinafter, each manufacturing condition will be described in detail.

  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, and if it exceeds 1300 ° C., the toughness deteriorates. Preferably, it is 1050-1250 degreeC.

Rolling end temperature: Ar3 temperature or higher. Ar3 temperature means the ferrite transformation start temperature during cooling, and can be determined by the following equation (c).
Ar3 = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo (c)
When the rolling end temperature is equal to or lower than the Ar3 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. Above the temperature.

  Immediately after completion of rolling, cooling is performed at a cooling rate of 5 ° C./s or more. When it is allowed to cool or gradually cool after the rolling is completed, precipitates are precipitated from the high temperature range, and the precipitates are easily coarsened and sufficient strengthening cannot be obtained. Therefore, it is an important production condition in the present invention to perform rapid cooling (accelerated cooling) to a temperature optimum for precipitation strengthening and prevent precipitation from a high temperature range. If the cooling rate is less than 5 ° C./s, the effect of preventing precipitation in a high temperature region is not sufficient and the strength is lowered. In addition, when the cooling start temperature is not more than the Ar3 temperature and ferrite is generated, the precipitation of fine precipitates is not obtained at the time of reheating, resulting in insufficient strength. Therefore, the cooling start temperature is set to the Ar3 temperature or more. About the cooling method at this time, it is possible to use arbitrary cooling equipment by a manufacturing process.

  Cooling stop temperature: 300 to 600 ° C. 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. If the cooling stop temperature is less than 300 ° C, bainite or martensite single phase structure may be formed, so that the HIC resistance is deteriorated. If it exceeds 600 ° C, ferrite transformation during reheating is not completed and pearlite is precipitated. Since the HIC resistance is deteriorated, the accelerated cooling stop temperature is defined as 300 to 600 ° C.

  Immediately after the accelerated cooling, reheating is performed to a temperature of 550 to 700 ° C. at a heating rate of 0.5 ° C./s or more. 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 with fine precipitates and soften the bainite phase and obtain a structure with a small strength difference between the ferrite phase and the bainite phase, reheat to a temperature range of 550 to 700 ° C. immediately after accelerated cooling. is required. In reheating, it is desirable to raise the temperature by at least 50 ° C. from the temperature after cooling. If the heating rate during reheating is less than 0.5 ° C./s, it takes a long time to reach the target reheating temperature, so that the production efficiency deteriorates and pearlite transformation occurs. Cannot be obtained and sufficient strength cannot be obtained. If the reheating temperature is less than 550 ° C, the ferrite transformation is not completed, and the untransformed austenite transforms to pearlite at the time of subsequent cooling, so that the HIC resistance deteriorates, and if it exceeds 700 ° C, the precipitate becomes coarse and sufficient strength is obtained. Therefore, the reheating temperature range is specified to be 550 to 700 ° C.

  There is no need to set the temperature holding time at the reheating temperature. If the production method of the present invention is used, even if it is cooled immediately after reheating, the ferrite transformation proceeds sufficiently, so that high strength due to fine precipitation can be obtained. In order to reliably complete the ferrite transformation, the temperature can be maintained for 30 minutes or less. However, if the temperature is maintained for more than 30 minutes, the precipitates may become coarse and the strength may be reduced. The cooling rate after reheating may be set as appropriate. However, since the ferrite transformation proceeds even in the cooling process after reheating, air cooling is preferable. As long as the ferrite transformation is not hindered, it is possible to perform cooling at a cooling rate faster than air cooling.

  As equipment for performing reheating up to a temperature of 550 to 700 ° C., a heating device can be installed on the downstream side of the cooling equipment for performing accelerated cooling. As the heating device, it is preferable to use a gas combustion furnace or induction heating device capable of rapid heating of the steel sheet. In addition, since the cooling rate after reheating may be arbitrary, it is not necessary to install special equipment in the downstream of a heating apparatus.

  In FIG. 2, the steel plate of the present invention (0.05 mass% C-0.15 mass% Si-1.25 mass% Mn-0.09 mass% Mo-0.01 mass% Ti) produced using the above production method is transmitted. The photograph observed with the scanning electron microscope (TEM) is shown. According to FIG. 2, it can be confirmed that very fine precipitates are deposited in a row, but this is due to transformation precipitation that causes precipitation at the austenite / ferrite interface during ferrite transformation. Extremely high precipitation strengthening is obtained. The precipitate is a carbide containing Mo and Ti, and this was confirmed by analysis using energy dispersive X-ray spectroscopy (EDX) or the like.

  In FIG. 3, the steel plate (0.05 mass% C-0.25 mass% Si-1.25 mass% Mn-0.04 mass% Nb-0.01 mass% Ti) of the present invention manufactured using the above manufacturing method is used. The photograph observed with the transmission electron microscope (TEM) is shown. According to FIG. 3, it can be confirmed that very fine precipitates are dispersed and precipitated, but this is due to transformation precipitation that causes precipitation at the austenite / ferrite interface during ferrite transformation, which is extremely high. Precipitation strengthening is obtained. The precipitate is a carbide containing Ti and Nb, and this was confirmed by analysis using energy dispersive X-ray spectroscopy (EDX) or the like.

  In FIG. 4, the schematic of an example of the manufacturing line for enforcing the manufacturing method of this invention is shown. As shown in FIG. 4, a hot rolling mill 3, an acceleration cooling device 4, an induction heating device 5, and a hot leveler 6 are arranged in the rolling line 1 from the upstream side toward the downstream side. By installing the induction heating device 5 or other heat treatment device on the same line as the hot rolling mill 3 that is a rolling facility and the accelerated cooling device 4 that is a subsequent cooling facility, reheating is performed quickly after the end of rolling and cooling. Since the treatment can be performed, the steel sheet after being rolled and accelerated and cooled can be immediately heated to 550 ° C. or higher.

  The steel plate of the present invention manufactured by the above manufacturing method is formed into a steel pipe by press bend forming, roll forming, UOE forming, etc., and transports crude oil or natural gas (electric-welded steel pipe, spiral steel pipe, UOE steel pipe) Etc. can be used. Since the steel pipe manufactured using the steel plate of the present invention has high strength and excellent HIC resistance, it is also suitable for transporting crude oil and natural gas containing hydrogen sulfide.

  Steel of chemical composition shown in Table 1 (steel types A to O) was made into a slab by a continuous casting method, and a thick steel plate (No. 1 to 27) having a plate thickness of 18 and 26 mm was produced using the slab.

  After the heated slab was rolled by hot rolling, it was immediately cooled using a water-cooled accelerated cooling facility and reheated using an induction heating furnace or a gas combustion furnace. The cooling equipment and induction heating furnace were in-line type. Table 2 shows the production conditions of each steel plate (No. 1 to 27).

  The microstructure of the steel sheet produced as described above was observed with an optical microscope and a transmission electron microscope (TEM). The components of the precipitate were analyzed by energy dispersive X-ray spectroscopy (EDX). The tensile properties and HIC resistance of each steel plate were measured. The measurement results are also shown in Table 2. Tensile properties were measured by performing a tensile test using a full thickness test piece in the rolling vertical direction as a tensile test piece, and measuring yield strength and tensile strength. In consideration of manufacturing variations, a steel having a yield strength of 480 MPa or more and a tensile strength of 580 MPa or more was evaluated as a high strength steel plate of API X65 grade or more. The HIC resistance is determined by performing an HIC test with an immersion time of 96 hours in accordance with NACE Standard TM-02-84. If no crack is observed, the HIC resistance is judged as good. Indicated. For the weld heat affected zone (HAZ) toughness, a Charpy test was performed using a test piece to which a heat history corresponding to a heat input of 40 kJ / cm was added by a reproducible thermal cycle apparatus. Then, those having Charpy absorbed energy (vE) at −30 ° C. of 70 J or more and ductile fracture surface ratio (SA) of 50% or more were considered good.

  In Table 2, Nos. 1 to 14 as examples of the present invention all have chemical components and production methods within the scope of the present invention, high yield strength of 480 MPa or more, tensile strength of 580 MPa or more, and HIC resistance. In addition, the weld heat-affected zone toughness was excellent. As for Ti, Mo, and some steel plates, fine carbide precipitates containing Nb and / or V and having a particle size of less than 10 nm were dispersed and precipitated. Moreover, the structure of the steel sheet was substantially a ferrite + bainite two-phase structure, and the fraction of the bainite phase was in the range of 10 to 80%. Further, fine TiN was dispersed and precipitated in the structure of the heat affected zone.

  Nos. 15 to 21 have chemical components within the scope of the present invention, but the manufacturing method is outside the scope of the present invention, so that the structure is not a ferrite + bainite two-phase structure, and fine carbides are dispersed. Since it did not precipitate, cracks occurred in the insufficient strength and in the HIC test. Nos. 22 to 27 have chemical components outside the scope of the present invention, so that coarse precipitates are not formed, and precipitates containing Ti and Mo are not dispersed and precipitated, so that sufficient strength cannot be obtained. Or cracks occurred in the HIC test, or the weld heat affected zone toughness was poor.

  In addition, when the reheating was performed in the induction heating furnace or in the gas combustion furnace, there was no particular difference in the results.

  Steel of chemical composition (steel types A2 to N2) shown in Table 3 was made into a slab by a continuous casting method, and a thick steel plate (No. 51 to 77) having a plate thickness of 18 and 26 mm was produced using this slab.

  After the heated slab was rolled by hot rolling, it was immediately cooled using a water-cooled accelerated cooling facility and reheated using an induction heating furnace or a gas combustion furnace. The cooling equipment and induction heating furnace were in-line type. Table 4 shows the production conditions of each steel plate (No. 51 to 77).

  The microstructure of the steel sheet produced as described above was observed with an optical microscope and a transmission electron microscope (TEM). The components of the precipitate were analyzed by energy dispersive X-ray spectroscopy (EDX). The tensile properties and HIC resistance of each steel plate were measured. The measurement results are also shown in Table 4. Tensile properties were measured by performing a tensile test using a full thickness test piece in the rolling vertical direction as a tensile test piece, and measuring yield strength and tensile strength. In consideration of manufacturing variations, a steel having a yield strength of 480 MPa or more and a tensile strength of 580 MPa or more was evaluated as a high strength steel plate of API X65 grade or more. The HIC resistance is determined by performing an HIC test with an immersion time of 96 hours in accordance with NACE Standard TM-02-84. If no crack is observed, the HIC resistance is judged as good. Indicated. For the weld heat affected zone (HAZ) toughness, a Charpy test was performed using a test piece to which a heat history corresponding to a heat input of 40 kJ / cm was added by a reproducible thermal cycle apparatus. And the Charpy absorbed energy (vE) at -30 ° C. was 70 J or more, and the ductile fracture surface ratio (SA) was 50% or more.

  In Table 4, the Nos. 51 to 64 as examples of the present invention all have chemical components and production methods within the scope of the present invention, high yield strength of 480 MPa or higher, tensile strength of 580 MPa or higher, and HIC resistance. In addition, the weld heat-affected zone toughness was excellent. The structure of the steel sheet was substantially a ferrite + bainite two-phase structure, and fine composite carbide precipitates containing Ti and Nb and / or V and having a particle size of less than 20 nm were dispersed and precipitated. Moreover, the structure of the steel sheet was substantially a ferrite + bainite two-phase structure, and the fraction of the bainite phase was in the range of 10 to 80%. Further, fine TiN was dispersed and precipitated in the structure of the heat affected zone.

  Nos. 65 to 71, the chemical components are within the scope of the present invention, but the manufacturing method is outside the scope of the present invention, so that the structure is not a ferrite + bainite two-phase structure, or a fine composite carbide. Was not dispersed and precipitated, resulting in insufficient strength and cracks in the HIC test. Nos. 72 to 77 have chemical components outside the scope of the present invention, so that coarse precipitates are generated and composite carbides containing Ti and Nb and / or V are not dispersed and precipitated. The strength was not obtained, cracks occurred in the HIC test, or the weld heat affected zone toughness was inferior.

  In addition, when the reheating was performed in the induction heating furnace or in the gas combustion furnace, there was no particular difference in the results.

The graph which shows the outline of the heat history in the manufacturing method of this invention. The photograph which observed the steel plate of the present invention with the transmission electron microscope (TEM). The photograph which observed the steel plate of the present invention with the transmission electron microscope (TEM). Schematic which shows an example of the manufacturing line for enforcing the manufacturing method of this invention.

Explanation of symbols

DESCRIPTION OF SYMBOLS 1 Rolling line 2 Steel plate 3 Hot rolling mill 4 Accelerated cooling device 5 Induction heating device 6 Hot leveler 10 Austenite single phase 11 Untransformed austenite 12 Bainite 13 Ferrite phase precipitation strengthened by fine precipitates 14 Tempered and softened bainite phase A Austenite region B Bainite region C Accelerated cooling D Reheating

Claims (6)

In mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 2%, P: 0.01% or less, S: 0.002% or less , Mo: 0.05 to 0.5%, Ti: 0.005 to 0.025%, Al: 0.07% or less, N: 0.004 to 0.006%, the balance being Fe and inevitable Consists of impurities , C / (Mo + Ti), which is the ratio of the amount of C in atomic% and the total amount of Mo and Ti, is 0.5 to 3, and the ratio of the amount of Ti and the amount of N in mass% A certain Ti / N is 2 to 4, the metal structure has a total volume fraction of ferrite and bainite of 95% or more , and precipitates containing Ti and Mo are dispersed and precipitated. High-strength steel sheet for line pipes with excellent HIC resistance and weld heat-affected zone toughness.   Furthermore, it contains Nb: 0.005 to 0.05% and / or V: 0.005 to 0.1% by mass%, and the total amount of C, Mo, Ti, Nb, and V in atomic% The ratio of C / (Mo + Ti + Nb + V) is 0.5 to 3, and composite precipitates containing Ti, Mo, Nb and / or V are dispersed and precipitated. A high-strength steel sheet for line pipes with excellent HIC resistance and weld heat-affected zone toughness as described in 1. In mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.5%, Mn: 0.5 to 2%, P: 0.01% or less, S: 0.002% or less Ti: 0.005 to 0.025%, Al: 0.07% or less, N: 0.004 to 0.006%, Nb: 0.005 to 0.05% and / or V: 0 0.005 to 0.10%, with the balance being Fe and inevitable impurities , and the ratio of the amount of C in atomic% and the total amount of Ti, Nb and V is C / (Ti + Nb + V) is 0.5 to 3 and Ti / N, which is the ratio of the Ti amount and the N amount in mass%, is 2 to 4, the total volume fraction of ferrite and bainite is 95% or more , and Ti and , Nb and / or V containing composite carbide is dispersed and precipitated, HIC resistance and weld heat affected zone toughness High-strength steel plate for line pipe superior in.   Furthermore, in mass%, Cu: 0.5% or less, Ni: 0.5% or less, Cr: 0.5% or less, Ca: 0.0005 to 0.003%, Mg: 0.005% or less The high-strength steel sheet for line pipes having excellent HIC resistance and weld heat-affected zone toughness according to any one of claims 1 to 3, characterized in that it contains one or more selected from the group consisting of:   The steel having the component composition according to any one of claims 1 to 4 is heated to a temperature of 1000 to 1300 ° C, hot-rolled at a rolling end temperature of Ar3 temperature or higher, and then 5 ° C / s or higher. HIC resistance resistance and weld heat affected zone characterized in that accelerated cooling is performed at a cooling rate of 300 to 600 ° C., and then reheating is performed immediately to 550 to 700 ° C. at a temperature rising rate of 0.5 ° C./s or more. A method for producing high-strength steel sheets for line pipes with excellent toughness.   A high-strength steel pipe excellent in HIC resistance and weld heat-affected zone toughness, characterized by being manufactured using the steel plate according to any one of claims 1 to 4.

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