JP4824142B2 - Steel for line pipe with good strength and ductility and method for producing the same - Google Patents

Description

The present invention relates to a high toughness, high strength, high ductility steel having a sufficient strength as a welded structural steel and having excellent ductility characteristics and excellent low temperature toughness, and a method for producing the same. In particular, the present invention relates to a steel for line pipes having strength and ductility excellent in elongation that requires low temperature toughness in a cold region and a method for producing the same.
This application claims priority on October 5, 2009 based on Japanese Patent Application No. 2009-231799 for which it applied to Japan, and uses the content here.

  In recent years, steels for line pipes are required to have higher strength in order to improve operational efficiency by improving safety, increasing the pressure of transport gas, and reducing costs by reducing the amount of steel used. In addition, the area where the steel material is used is expanding to an area where the natural environment such as a cold region is severe, and severe toughness characteristics are required. In particular, structural steels used in earthquake-prone areas are required to have plastic deformability and ductile fracture characteristics in addition to the conventional required characteristics.

  For example, Patent Document 1 presents steel aimed at obtaining a high uniform elongation in order to suppress ductile fracture. Here, a mixed structure in which an appropriate amount of a hardened phase is introduced into ferrite is mixed by quenching, two-phase treatment, and tempering treatment (QLT treatment) to achieve high ductility. Moreover, in patent document 2, high ductility is aimed at by optimization and accelerated cooling of a steel component and quenching hardenability (Di).

Patent Document 3 presents a steel sheet having excellent HIC resistance. A steel sheet of 16 mm or less and X56 or less is manufactured without acceleration cooling, and a mixed structure composed of ferrite-pearlite-bainite is manufactured to ensure HIC resistance.
Patent Document 4 presents a steel plate of 570 MPa or more. After accelerated cooling, the generation of island martensite (MA) is suppressed by induction heating, and the hardness of the surface layer is suppressed. High strength and toughness are achieved by suppressing variations in hardness in the thickness direction.

  In general, high strength steel is required to increase the carbon equivalent and quenching index in order to obtain strength. However, when the carbon equivalent is simply increased, ductility and toughness are reduced. On the other hand, steel for large-diameter line pipes is required to reduce variations in strength and ductility within the plate in order to manage ductility after pipe making, such as UOE and JCOE. It is necessary to suppress this.

JP 2003-253331 A JP 2003-288512 A JP 2001-158936 A JP 2008-121036 A

The present inventor has found that the prior art has the following problems.
Steel for large diameter line pipes is required to reduce variations in properties such as strength and ductility in the plate in order to manage ductility after pipe making such as UOE and JCOE. Therefore, for example, a technique for reducing the in-plate variation in steel characteristics by forming a uniform structure by QLT (Quenching-Lamellarizing-Tempering) processing is adopted. However, the QLT process is expensive because it is heat-treated at least three times at a high temperature. Further, it is possible to achieve high strength and high ductility by accelerated cooling corresponding to the two-phase region heat treatment, but it is extremely difficult to make the cooling in the plate uniform in accelerated cooling.

  In order to ensure the HIC resistance as described in Patent Document 3, it is required to reduce the hardness of the steel as a whole by Vickers hardness of 248 Hv or less as specified by NACE. For this reason, it is difficult to increase the strength. Moreover, when such steel is manufactured by accelerated cooling, the production | generation of bainite and MA which are hardened structures is inevitable. In this case, a decrease in elongation (particularly local elongation) is inevitable.

  As described in Patent Document 4, after the accelerated cooling, the local elongation can be improved by reducing the hardness of the surface layer by induction heating. However, a band structure is not formed in a structure manufactured by accelerated cooling and called a mixed structure of bainite and ferrite, as shown in FIG. 2 of the publication, and general ferrite is not recognized. In such a structure, the local elongation is improved, but the uniform elongation is significantly reduced, and the total elongation is reduced. In addition, it is expensive to install induction heating equipment in the production line, and heat treatment by induction heating is extremely difficult to control because of the depth of the skin due to electromagnetic waves.

  Accordingly, the present invention provides an inexpensive high-strength thick steel plate having good toughness and ductility characteristics in a steel for line pipes, a manufacturing method thereof, and a quality control method thereof.

  As a result of investigating the influence of hardened structure such as MA on the ductility and segregation that promotes the formation of MA, the present inventors have made the steel structure a two-phase structure of ferrite and pearlite from the viewpoint of strength and ductility balance. And it discovered that the fall of local elongation was prevented by suppressing the production | generation of MA.

The gist of the present invention is as follows.
(1) The steel for line pipes with good strength and ductility according to one embodiment of the present invention is mass%, C: 0.07 to 0.15%, Si: 0.05 to 0.60%, Mn: 0.80 to 1.80%, P: 0.010 or less, S: 0.007% or less, V: 0.05 to 0.12%, Nb: 0.005 to 0.070%, Al: 0. 005 to 0.08%, Ti: 0.005 to 0.030%, Ca: 0.0005 to 0.0035%, N: 0.0020 to 0.0060%, O: 0.0030% or less The balance is composed of iron and inevitable impurities, the structure is a two-phase structure of ferrite and pearlite, the area fraction of island martensite is less than 1.5%, the plate thickness is 18 mm or more, and the yield strength Is 450 MPa or more.
(2) The steel for line pipes of the above (1) is further in mass%, Cu: 0.05 to 0.70%, Ni: 0.05 to 0.70%, Cr: 0.80% or less, Mo: 0.30% or less, B: 0.0003 to 0.0030%, Mg: 0.0003 to 0.0030%, REM: 0.0005 to 0.0050% Also good.
(3) In the steel for line pipes according to (1) or (2), the bainite structure may not be detected under an optical microscope.
(4) In the steel for line pipes of (1) or (2) above, the segregation degree of Mn may be 1.7 or less.
(5) In the steel for line pipes of (4) above, the segregation degrees of Si and P may be 1.5 or less and 8.0 or less, respectively.
(6) According to another aspect of the present invention, a method for producing a steel for a line pipe with good strength and ductility is obtained by adding a slab having a chemical component of the steel for a line pipe according to (1) or (2) to 1250 ° C. or less. After heating to a temperature of 850 ° C., hot rolling is performed at a cumulative reduction ratio of 40% or higher in a temperature range of 850 ° C. or higher, and hot rolling is terminated in a temperature range of 700 to 800 ° C., followed by air cooling.
(7) In the manufacturing method of (6), after the air cooling, the steel sheet may be tempered in a temperature range of 500 to 300 ° C.
(8) A quality control method for steel for line pipes with good strength and ductility according to another aspect of the present invention is the quality control method for steel for line pipes according to (1) or (2) above. The ductility is improved by measuring the amount of island martensite in the steel sheet produced by hot rolling and controlling the production conditions so that the area fraction of island martensite is less than 1.5%.
(9) In the quality control method of (8) above, the ductility may be improved by setting the cooling method after hot rolling to air cooling.
(10) In the quality control method of (8) or (9) above, the ductility may be improved by setting the segregation degree of Mn at the steel plate stage to 1.7 or less.
(11) In the quality control method of the above (8) or (9), the ductility may be improved by setting the bainite structure fraction at the steel sheet stage to 1% or less.

  As described above, according to the present invention, a steel for line pipes having good strength and ductility can be obtained, which is extremely useful industrially.

It is a graph which shows the relationship between MA, uniform elongation (U.El), local elongation (L.El), and total elongation (T.El).

  In general, a large amount of alloy addition and accelerated cooling are effective for increasing the strength, but the ductility is deteriorated on the contrary because it becomes a structure with high hardenability. In addition, when the strength is increased, MA or baked bainite may be locally generated, but the relationship between these and elongation is unclear.

Therefore, the present inventors have conducted detailed studies on the influence of the structure on the ductility, investigated the influence of the hardened structure such as MA on the ductility and the segregation that promotes the formation of MA, and the following is necessary. It revealed that.
1. From the viewpoint of strength and ductility balance, it is necessary to have a two-phase structure of ferrite and pearlite.
2. The production of MA has little influence on the uniform elongation, but as the MA fraction increases, the local elongation deteriorates significantly. Also, bainite of about 400 Hv or more shows the same behavior as MA, and its generation reduces local elongation.
3. In order to suppress the formation of MA, it is important to control the manufacturing process, but it is also important to reduce segregation. In particular, as for segregation, when the degree of segregation increases, band-like MA is formed and the local elongation is remarkably reduced.

As described above, the ductility of a material which is generally increased in strength for line pipes is low. For example, when a bainite single phase structure is formed by using accelerated cooling, it is easy to secure a strength of about 450 MPa class in terms of YS (yield strength). However, the ductility (uniform elongation) decreases with increasing strength. In the case of a single-phase structure, particularly with respect to ductility, local elongation is remarkably reduced, and it is difficult to ensure a balance between strength and ductility. Further, when a ferrite single phase is used, it is possible to increase ductility, but it is difficult to ensure strength. Further, if there is ferrite and a hardened bainitic structure generated by accelerated cooling, local elongation is lowered.
For this reason, a two-phase structure of ferrite for ensuring high ductility and pearlite for ensuring strength is required.

  Moreover, when strengthening, it is inevitable that MA is generated in part of pearlite or bainite. An example showing the relationship between this MA and uniform elongation (U. El), local elongation (L. El) and total elongation (T. El) is shown in FIG. As shown in FIG. 1, due to the increase in MA, the uniform elongation hardly decreases but the local elongation decreases. It can be seen that the overall elongation decreases as a result of the decrease in local elongation due to the increase in MA. In order to avoid a decrease in uniform elongation, the MA is limited to 1.5% or less in the present invention. More preferably, MA is limited to 1.0% or less or 0.5% or less.

  Therefore, in the present invention, the ductility is ensured in spite of the high strength steel by controlling the two-phase structure of ferrite and pearlite and controlling the MA and segregation at low cost.

  The configuration of the present invention will be described in detail below. First, the reasons for limiting the composition of the steel material of the present invention will be described. In the present specification, the content% means mass%.

C: 0.07 to 0.15%
C is an element necessary for ensuring strength, and it is necessary to add 0.07% or more. Since a large amount of addition may cause a decrease in ductility and low temperature toughness, the upper limit is set to 0.15%. Desirably, 0.12% or less is good.

Si: 0.05-0.60%
Si is an element effective as a deoxidizing element and increases the strength of the steel by solid solution strengthening. However, when its content is less than 0.05%, these effects are not recognized. Further, if added over 0.60%, a large amount of MA is generated in the structure, so that the toughness is deteriorated. Therefore, the addition amount of Si is set to 0.20 to 0.60%. Note that since MA (or retained austenite) starts to increase at 0.45% or more, it is preferably less than 0.45%.

Mn: 0.80 to 1.80%
Mn is an effective element for increasing the strength because it increases the strength of the steel. For that purpose, addition of 0.80% or more is necessary. However, if it exceeds 1.80%, the degree of segregation such as center segregation and microsegregation increases, so that the formation of MA is promoted and the local elongation is lowered. For this reason, the appropriate range of the amount of Mn added is set to 0.80 to 1.80%. In order to increase the strength, it is desirable that the lower limit of Mn is 1.0%, 1.2% or 1.3%.

P: 0.010% or less When P exceeds 0.010%, it segregates at the grain boundaries and significantly deteriorates the toughness of the steel. In addition, the concentration of P in the segregation zone is increased to promote the formation of MA. For this reason, the upper limit of the addition amount was set to 0.010%. In addition, it is desirable to reduce as much as possible from a viewpoint of suppressing the fall of a ductility and a toughness value.

S: 0.007% or less S is present in steel by forming MnS, and has the effect of refining the structure after rolling and cooling. However, if it exceeds 0.007%, the toughness of the base metal and the welded portion is reduced. Deteriorate. For this reason, S was made into 0.007% or less.

Nb: 0.005 to 0.070%
In the present invention, since the cooling is air cooling for the structure control, Nb is an important element for securing the strength. Moreover, by adding Nb, the heated austenite at the time of slab reheating or quenching becomes finer and the strength of the steel is increased. Therefore, it is necessary to add 0.005% or more. However, since excessive Nb addition reduces the ductility of the base material, the upper limit of the Nb addition amount is set to 0.070%. In order to improve the toughness of the base material, the upper limit of Nb may be limited to 0.050% or 0.35%.

V: 0.05-0.12%
V has almost the same function as Nb, but its effect is smaller than that of Nb. The effect similar to Nb is less effective at less than 0.05%. However, if it exceeds 0.12%, the ductility deteriorates. For this reason, the appropriate range of the addition amount of V was made 0.05 to 0.12%.

Al: 0.005 to 0.08%
For the purpose of deoxidation, Al needs to be added in an amount of 0.005% or more. If it is less than 0.005%, the formation of MA is suppressed, but it becomes weak deoxidation and the generation probability of coarse oxides is increased, so the local elongation is lowered. On the other hand, excessive addition exceeding 0.08% reduces weldability. This problem is particularly remarkable in SAW (Submerged Arc Welding) using flux, etc., which deteriorates the toughness of the weld metal and decreases the HAZ (Heart Affected Zone) toughness. For this reason, the upper limit of Al was made 0.08%. In order to improve weldability, it is preferable to limit the upper limit of Al to 0.06 or 0.04%.

Ti: 0.005-0.030%
Addition of 0.005% or more is desirable for Ti to combine with N to form TiN effective for high strength and high ductility in steel. However, if Ti is added in an amount exceeding 0.030%, TiN may be coarsened and the ductility of the base material may be reduced. For this reason, Ti was taken as 0.005 to 0.030% of range.
Ca: 0.0005 to 0.0035%,
When Ca is added in an amount of 0.0005% or more, it has an effect of controlling the form of sulfide (MnS), increasing the absorbed energy of Charpy and improving the low temperature toughness. However, if it exceeds 0.0035%, a large amount of coarse CaO and CaS is generated, which adversely affects the ductility and toughness of the steel. For this reason, 0.0035% was limited to the upper limit of the Ca content.

N: 0.0020 to 0.0060%
N needs to be added in an amount of 0.0020% or more in order to form TiN effective in increasing the strength and ductility in steel by bonding with Ti. However, since N has a very large effect as a solid solution strengthening element, adding a large amount may cause deterioration of ductility. Therefore, the upper limit of N is set to 0.0060% so that the effect of TiN can be maximized without greatly affecting the ductility.

O: 0.0030% or less In the case of high-strength steel, if O exceeds 0.0030%, the cleanliness and toughness of the steel are deteriorated. Therefore, the upper limit value is set to 0.0030%.

  The basic components in the present invention are as described above and can sufficiently achieve the target value, but in order to further improve the characteristics, one or more of the following elements are added as selective elements as necessary. Also good.

Cu: 0.05 to 0.70%
Cu is an element advantageous for increasing the strength. In order to ensure the precipitation effect by Cu, addition of 0.05% or more is desirable. However, excessive addition increases the hardness of the base material and decreases the ductility, so the upper limit was made 0.70%. In order to improve ductility, it is preferable to limit Cu to 0.50% or less, 0.30% or less, or 0.20% or less.

Ni: 0.05-0.70%
Ni improves strength and toughness without adversely affecting weldability and the like, and is also effective in preventing Cu cracking. In order to obtain these effects, addition of 0.05% or more is necessary. However, since Ni is expensive, if it is made 0.70% or more, it becomes impossible to manufacture steel at a low cost, so the addition amount is limited to 0.70% or less. In order to improve ductility, it is desirable to add less Ni, and the content may be limited to 0.50% or less, 0.30% or less, or 0.20% or less.

Cr: 0.80% or less Cr is an element that increases the strength of the base material. However, if it exceeds 0.80%, the hardness of the base material is increased and the ductility is deteriorated. Therefore, the upper limit is set to 0.80%. In order to improve ductility, it is preferable to limit Cr to 0.50% or less, 0.30% or less, or 0.10% or less.

Mo: 0.30% or less Cr, like Mo, is an element that increases the strength of the base material. However, if it exceeds 0.30%, the hardness of the base material is increased and the ductility is deteriorated. Therefore, the upper limit is set to 0.30%. In order to improve ductility, it is preferable to limit Mo to 0.20% or less or 0.10% or less.

B: 0.0003 to 0.0030%
B is an element that dissolves in steel to increase the hardenability and increase the strength. In order to obtain this effect, addition of 0.0003% or more is necessary. However, if B is added excessively, the base material toughness is lowered. For this reason, the upper limit was made 0.0030%.

Mg: 0.0003 to 0.0030%
Mg also suppresses the growth of austenite grains, has the effect of maintaining fine grains, and improves toughness. In order to enjoy this effect, addition of at least 0.0003% or more is essential, and this amount is set as the lower limit. On the other hand, if the amount added is increased more than necessary, not only does the effect margin for the amount added become small, but Mg does not necessarily have a high steelmaking yield, so the economy is lost. In view of these, the upper limit is limited to 0.0030% in the present invention.

REM: 0.0005 to 0.0050%
REM, like Mg, also suppresses the growth of austenite grains, keeps them fine, and improves toughness. In order to enjoy this effect, at least 0.0005% or more must be added, and this amount is set as the lower limit. On the other hand, if the amount added is increased more than necessary, not only does the effect margin for the amount added become small, but Mg does not necessarily have a high steelmaking yield, so the economy is lost. In view of these, the upper limit is limited to 0.0050% in the present invention.

  In the present invention, it is possible to manufacture UOE and JCOE steel pipes having high strength and high ductility mainly as steel materials for line pipe welding. One feature of the steel according to the present invention is that the composite properties of strength, toughness and ductility required for the line pipe are ensured mainly by a two-phase structure of ferrite (having a ferrite band structure) and pearlite. The ferrite structure in the present invention has a so-called ferrite band structure along the hot rolling direction in the structure.

Further, at this time, if the MA fraction is 1.5% or more, a large amount of voids are generated in the vicinity of the MA during the tensile test, and the shear elongation is promoted by the plastic flow, and the local elongation is remarkably reduced. In order to suppress the segregation-induced MA that deteriorates the local elongation and to make the MA less than 1.5%, it is important not to cool rapidly. More specifically, MA having a Vickers hardness of 400 to 700 Hv causes the generation of voids particularly frequently and causes a significant deterioration in local elongation. For this reason, if 400-700Hv MA is suppressed, degradation of local elongation can be avoided. As long as the MA fraction is in the above range, the segregation degree of each element in the steel sheet is not necessarily regulated in the present invention. However, the segregation degree of the Mn steel plate is preferably 1.7 or less. More preferably, the segregation degrees of the Si and P steel sheets are 1.5 and 8.0 or less, respectively.
When the Mn segregation degree of the steel sheet exceeds 1.7, the Si segregation degree exceeds 1.5, or the P segregation degree exceeds 8.0, the formation of MA becomes significant. In the present invention, only the element segregation degree of the steel sheet after rolling is specified, but in order to ensure that the Mn segregation degree is 1.7 or less when the sheet is formed, the Mn segregation degree in the slab stage is 1.1 or less. There is a need. Moreover, in this invention, in order to control Mn segregation of the final steel plate, the manufacturing method of a slab is not limited. However, in order to control the MA distribution, it is necessary to reduce not only macro segregation such as center segregation but also micro segregation.

In the present invention, the degree of segregation means component analysis in a 1 mm ² region in the plate thickness direction 1/2, and the peak concentration of Mn, Si or P at these positions divided by the average concentration of each element. Say. Component analysis can use, for example, EPMA (Electron Probe Micro Analyzer) or CMA (Computer-aided Micro Analyzer).

  Next, the reason for limiting the production conditions of the steel of the present invention will be described.

  The Mn segregation degree of the slab is preferably 1.1 or more. If a steel plate is manufactured from such a slab, the segregation degree of Mn of a steel plate can be reliably made 1.7 or less. As described above, when it exceeds 1.7, the generation of MA tends to be remarkable. In the present invention, only the segregation degree of Mn is defined for the slab. Although segregation of Si and P is important, it is not specified because the influence of the manufacturing process is larger than that of Mn. If allowed by the manufacturing process, the segregation of Si and P is desirably 1.5 or 8.0 or less. As a method for controlling the Mn segregation degree of the slab, for example, a widely known method such as light reduction (soft reduction), electromagnetic stirring during continuous casting, diffusion heat treatment of the segregation element by high-temperature heat treatment to the slab, and the like should be used. Can do. In the manufacturing process of the slab used in the present invention, light reduction was performed.

When the heating temperature exceeds 1250 ° C., the crystal grain size becomes very coarse, and a large amount of scale is generated on the steel surface, and the quality of the surface is significantly lowered. For this reason, the upper limit of reheating temperature was 1250 degreeC.
It is necessary to perform hot rolling with a cumulative rolling reduction of 40% or more in a temperature range of 850 ° C. or higher. The increase in the amount of reduction in this temperature range contributes to the refinement of the austenite grains during rolling, and as a result, has the effect of refining the ferrite grains and improving the mechanical properties. In order to obtain such an effect, the cumulative rolling reduction needs to be 40% or more in the temperature range of 850 ° C. or higher.

  It is necessary to perform hot rolling with a cumulative rolling reduction of 40% or more in the non-recrystallization temperature range. The increase in the amount of reduction in the non-recrystallization temperature range contributes to the refinement of austenite grains during rolling, and as a result, has the effect of refining ferrite grains and improving mechanical properties. In order to obtain such an effect, the cumulative rolling reduction in the non-recrystallized region needs to be 40% or more. For this reason, the cumulative reduction amount in the non-recrystallized region is limited to 40% or more.

  The steel slab needs to be air-cooled after hot rolling is completed in a temperature range of 800 to 700 ° C. In this case, it is desirable to cool slowly at a cooling rate of less than 5 ° C./s. In the present invention, rolling is completed at a two-phase temperature of 800 to 700 ° C., and a mixed structure of ferrite and pearlite is generated. Thereby, base material toughness such as DWTT, high strength, and high ductility are achieved. Although depending on the chemical component, it is more desirable that the hot rolling is in a temperature range of 780 to 720 ° C. from the viewpoint of securing a balance of strength, ductility, and toughness.

When the rolling end temperature exceeds 800 ° C., a band-like pearlite structure is not formed, and ductility and base material toughness are lowered. Moreover, when it becomes less than 700 degreeC, the amount of work ferrite will increase and ductility (uniform elongation) will fall remarkably.
In the present invention, the cooling method is defined only as air cooling, and the cooling rate is determined in relation to the plate thickness. With a general plate thickness, when the cooling rate exceeds 5 ° C./s, MA and bainite are likely to be formed, and the toughness and ductility are reduced. For this reason, slow cooling at a cooling rate of less than 5 ° C./s is desirable. More preferably, it is 2 ° C./s or less. By performing air cooling, the above cooling rate can be easily obtained.

  It is effective to perform a warm tempering treatment in the temperature range of 500 to 300 ° C. after cooling. If the tempering temperature exceeds 500 ° C., the strength is lowered, and if it is less than 300 ° C., the MA is not decomposed, so that the ductility is lowered. Further, the warm tempering treatment in the temperature range of 500 to 300 ° C. improves local elongation from the viewpoint of dehydrogenation.

  In the present invention, the structure needs to be a two-phase structure of ferrite and pearlite as described above. The structure fraction is not necessarily specified, but preferably the ferrite fraction is about 60 to 95%.

  As described above, when a large amount of hardened tissue enters, local elongation is reduced. For this reason, in this invention, the production | generation of bainite is avoided. In the steel sheet produced under the conditions of the present invention, bainite is not detected when measured with an optical microscope. As a result, the present invention steel has a two-phase structure of ferrite and pearlite. However, industrially, it is difficult to completely remove bainite, and in the present invention, a bainite structure is confirmed at the electron microscope level. As the tissue fraction, bainite is desirably 1% or less.

  In the present invention, the plate thickness is 18 mm or more, and the yield strength is 450 MPa or more. When the plate thickness is less than 18 mm, it is easy to ensure the yield strength. However, even with air cooling, the cooling rate increases, so that a large amount of a hardened structure such as bainite is generated and it is difficult to ensure the elongation. In general, it is known that the elongation improves as the yield strength decreases. If it is less than 540 MPa, it is possible to easily obtain a total elongation of 20% or more in the GOST tensile test and 40% or more in the API tensile test without controlling the chemical composition, structure, manufacturing method and the like. For this reason, this invention exhibits an especially suitable effect in the steel plate whose yield strength is 450 Mpa or more. More preferably, the steel sheet may have a yield strength of 540 MPa or more. The toughness of the steel sheet is preferably 70% or more at a ductile fracture surface ratio of −20 ° C. according to the DWTT test.

  Next, examples of the present invention will be described.

A slab obtained by casting molten steel having the chemical components shown in Table 1 was hot-rolled under the conditions shown in Table 2 to obtain a steel plate. In Table 2, the “cumulative reduction amount” column indicates the cumulative reduction rate in a temperature range of 850 ° C. or higher.
This was followed by a test to evaluate the mechanical properties. In the cooling method column of Table 2, the comparative steels r and s were accelerated and cooled at the rate indicated in parentheses (° C./s). The steel plates other than the above were cooled by air cooling.

  For tensile test specimens, Russian standard GOST test specimens were taken from each sample steel plate, and YS (yield strength: 0.5% underload), TS (tensile strength), and total elongation (total elongation: T. El). The uniform elongation (U. El) and the local elongation (L. El) were evaluated. As for the base material toughness, a ductile fracture surface ratio of −20 ° C. was evaluated by a DWTT test, and 70% or more was determined to be acceptable.

  Table 3 summarizes the mechanical properties of each steel. Steels a to o are examples of the present invention. As is apparent from Tables 1 and 2, these steel sheets satisfy the requirements of chemical components and production conditions, and as shown in Table 3, the base metal strength, ductility and toughness are good. The steel structure was a ferrite + pearlite two-phase structure (under an optical microscope), and MA was 1.5% or less.

  On the other hand, steels p to ad are comparative steels that depart from the scope of the present invention. The steel p to v and the ab to ad have different production conditions from the present invention, and the chemical components of the steel w to aa are outside the scope of the present invention. As a result, the steels p to ad were inferior to the steel of the present invention in one or more of the mechanical properties of the base material.

表３中の「組織」欄は鋼板中の組織を示し、Ｆがフェライト、Ｐがパーライト、Ｂがベイナイトを表す。

The “structure” column in Table 3 indicates the structure in the steel sheet, where F is ferrite, P is pearlite, and B is bainite.

As can be seen with reference to Tables 1 to 3, the toughness of steel p decreased because of its low cumulative reduction. Steel q had a high rolling end temperature and was unable to control the structure, so the toughness decreased. Since the steel r and steel s were accelerated and cooled, a large amount of bainite and MA were produced in the process, resulting in a decrease in ductility. In steel r, MA is decomposed by tempering, but because there is bainite, the toughness is recovered, but the recovery of elongation is insufficient. Steels t, ab to ad intentionally produced a large amount of MA by increasing the segregation degree of the plate by omitting the light reduction treatment in the slab manufacturing stage. For this reason, the elongation is reduced. Since steel u has a small plate thickness and a high cooling rate, MA is primarily generated before tempering. Steel u is subsequently tempered, but since the temperature is low, MA is not decomposed and elongation and toughness cannot be obtained. Steel v has a high tempering treatment temperature, and the amount of MA is reduced, but the yield strength is low.
Since the steel C has a low C content, the strength of the base material has decreased. Moreover, since steel x has a high Si content, MA increased and ductility decreased. Since the steel y has a high Mn content, the increase in MA and the MA itself harden, so that predetermined elongation characteristics and toughness cannot be obtained. Steel z is weakly deoxidized because of its low Al content. Steel aa has a high Ca content. For this reason, a relatively coarse oxide is produced in steels z and aa, and sufficient elongation cannot be obtained.

  Moreover, although the slabs 1-7 which satisfy the composition of the present invention are used in the comparative steels p to v, since the manufacturing conditions are different from those of the present invention, the elongation and / or toughness is inferior.

  As described above, according to the present invention, steel for line pipes having good strength and ductility can be obtained, which is extremely useful industrially.

Claims (11)

% By mass
C: 0.07 to 0.15%,
Si: 0.05 to 0.60%,
Mn: 0.80 to 1.80%,
P: 0.010 or less,
S: 0.007% or less,
V: 0.05 to 0.12%,
Nb: 0.005 to 0.070%,
Al: 0.005 to 0.08%,
Ti: 0.005 to 0.030%,
Ca: 0.0005 to 0.0035%,
N: 0.0020 to 0.0060%,
O: contains 0.0030% or less, the balance consists of iron and inevitable impurities,
The structure is a two-phase structure of ferrite and pearlite,
The area fraction of island martensite is less than 1.5%,
A steel for line pipes with good strength and ductility, characterized in that the plate thickness is 18 mm or more and the yield strength is 450 MPa or more. Furthermore, in mass%,
Cu: 0.05 to 0.70%,
Ni: 0.05 to 0.70%,
Cr: 0.80% or less,
Mo: 0.30% or less,
B: 0.0003 to 0.0030%,
Mg: 0.0003 to 0.0030%,
REM: 0.0005-0.0050% of 1 type, or 2 or more types were contained, Steel for line pipes with favorable intensity | strength and ductility of Claim 1 characterized by the above-mentioned.   The steel for line pipes having good strength and ductility according to claim 1 or 2, wherein a bainite structure is not detected under an optical microscope.   The steel for line pipes having good strength and ductility according to claim 1 or 2, wherein the segregation degree of Mn is 1.7 or less.   The steel for line pipes with good strength and ductility according to claim 4, wherein the segregation degrees of Si and P are 1.5 or less and 8.0 or less, respectively.   The slab of the chemical component according to claim 1 or 2 is heated to a temperature of 1250 ° C or lower, and then hot-rolled at a cumulative reduction ratio of 40% or higher in a temperature range of 850 ° C or higher, and a temperature of 700 to 800 ° C. A method for producing steel for line pipes with good strength and ductility, characterized by air cooling after finishing hot rolling in a region.   The method for producing steel for line pipes with good strength and ductility according to claim 6, wherein after the air cooling, the steel plate is tempered in a temperature range of 500 to 300 ° C. A quality control method for steel for line pipes according to claim 1 or 2,
Measure the amount of island martensite in the steel sheet manufactured by hot rolling,
A quality control method for steel for line pipes having good strength and ductility, wherein ductility is improved by controlling production conditions so that the area fraction of island martensite is less than 1.5%.   The quality control method for steel for line pipes with good strength and ductility according to claim 8, wherein ductility is improved by air cooling as a cooling method after hot rolling.   The quality control method for steel for line pipes with good strength and ductility according to claim 8 or 9, wherein the ductility is improved by setting the segregation degree of Mn at a steel plate stage to 1.7 or less.   The quality control method for steel for line pipes with good strength and ductility according to claim 8 or 9, wherein the ductility is improved by setting the bainite structure fraction in the steel sheet stage to 1% or less.

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