Patent Application
20080047635
The present invention relates to a heavy wall seamless steel pipe for line pipe excellent in strength, toughness and weldability, and a manufacturing method thereof. The heavy wall seamless steel pipe means a seamless steel pipe having a wall thickness of 25 mm or more. The seamless steel pipe of the present invention is a high-strength, high-toughness heavy wall seamless steel pipe for line pipe having a strength of not less than X70 regulated in API (American Petroleum Institute) Standard, that is, a strength of X70 (yield strength of 482 MPa or more), X80 (yield strength of 551 MPa or more), X90 (yield strength of 620 MPa or more), X100 (yield strength of 689 MPa or more), and X120 (yield strength of 827 MPa or more), which is particularly suitably used for submarine flow lines.
In recent years, petroleum and gas resources of oil fields located on land or in shallow sea areas are running dry, and development of offshore oil fields in deep sea areas has been increasingly activated. In deep-sea oil fields, there is a need for the transportation of crude oil or gas from a pit of an oil well or gas well set on the sea bottom to a platform on the ocean by the use of a flow line or riser.
An inner part of a pipe constituting the flow line laid in the deep sea suffers a high internal fluid pressure in addition to a deep stratum pressure, and is also subjected to repeated strains by ocean waves and influenced by the sea water pressure of the deep sea at the time of shutdown. Therefore, a heavy wall steel pipe with high strength and toughness is desired as the pipe used for this purpose.
Such a seamless steel pipe with high strength and toughness has been manufactured by piercing a billet heated to high temperature by a piercing mill, shaping into a pipe shape of product by rolling and drawing, and then performing heat treatment. In recent years, however, simplification of the manufacturing process by applying an in-line heat treatment has been examined from the viewpoint of energy and process saving. Particularly, paying attention to effective use of the heat possessed by the material after hot-worked, a process for performing quenching without once cooling beforehand to room temperature has been introduced. According to this method, substantial energy saving and increased efficiency of the manufacturing process can be attained, enabling significant reduction in manufacturing cost.
A steel pipe manufactured in the in-line heat treatment process of performing quenching directly after finish rolling has not been subjected to transformation and reverse transformation since, unlike in the past, it is not reheated after once cooling to room temperature and rolling. Therefore, the grains are apt to be coarsened, and it is not easy to ensure the toughness and corrosion resistance. Some techniques have been proposed in order to make fine grains of a finish-rolled steel pipe and ensure the toughness or corrosion resistance even if the grains are not so fine.
For example, the following Patent Document 1 (Japanese Patent Unexamined Publication 2001-240913) discloses a technique for making fine grains by adjusting the leading time to let it into the reheating furnace after finish rolling. The following Patent Document 2 (Japanese Patent Unexamined Publication 2000-104117) discloses a technique for adjusting the chemical composition, particularly, the contents of Ti and S to provide a satisfactory performance even with a relatively large grain size.
However, the technique disclosed in Patent Document 1 cannot respond to manufacture of a heavy wall steel pipe with high strength for offshore oil fields in depth, which has been increasingly demanded in recent years. For example, the heavy wall steel pipe requires a high finish rolling temperature, and it takes an excessive time to ensure an intended reheating furnace temperature, and seriously reduces the production efficiency. The method described in Patent Document 2 is also hardly applicable to heavy wall materials. Since the cooling rate in the in-line heat treatment is reduced in the case of heavy wall materials, the toughness is deteriorated even if steel of the composition disclosed in Patent Document 2 is applied.
[Patent Document 1] Japanese Patent Unexamined Publication 2001-240913.
[Patent Document 2] Japanese Patent Unexamined Publication 2000-104117.
In order to solve the above-mentioned problems, it is the objective of the present invention to provide a seamless steel pipe for line pipe having high strength and stable toughness, particularly, in a steel pipe with a heavy wall and also a manufacturing method thereof.
1. Fundamental Examination and Findings
Factors governing the toughness of a heavy wall seamless steel pipe were analyzed first. As a result, the following information was found.
(1) Cooling conditions in and after solidification of molten steel greatly influence the toughness. Since a lower cooling rate causes reduction in toughness, the cooling must be accomplished at not less than a specific cooling rate.
(2) A blooming process for heating an ingot to a high-temperature range for hot working does not have a good influence on the toughness.
(3) The above-mentioned reduction in toughness is caused by the precipitation form of Ti carbonitride due to the cooling rates in and after solidification. In order to prevent this reduction in toughness, it is important to finely precipitate the Ti carbonitride.
(4) Precipitation strengthening deteriorates the balance between strength and toughness in the case of in-line heat treatment materials. Although it is disadvantageous for obtaining a high strength, it is desirable to utilize the transformation strengthening and the solid-solution strengthening, without utilizing the precipitation strengthening, in order to obtain a high toughness.
(5) It is necessary to prevent from the generation of a retained austenite and a low-temperature transformed martensite in order to obtain a homogenous metal microstructure.
(6) Regarding the chemical composition, it is desirable to reduce the contents of Si, P and S, and to control the contents of Nb and V so as not to exceed the specific upper limits, and also to include a proper quantity of Ti, in addition to a proper quantity of at least one elements selected from Ca, Mg and REM. Accordingly, the toughness of heavy wall materials will be significantly improved.
(7) The findings described in (1) to (6) above were obtained on the assumption of the in-line heat treatment. However, if it is applied to a steel pipe subjected to an off-line heat treatment, an increased toughness can be obtained. Therefore, the above-mentioned findings can also be used to manufacture a high-strength material by an off-line heat treatment.
2. Basic Test and Result
Since an in-line heat treatment does not have a fine-grain making process by “transformation-reverse transformation” unlike an off-line heat treatment, a fine-graining itself at the end of rolling is required to ensure the toughness.
It is generally said that although the as-solidified grains are coarse, the grains become fine by reheating in order to perform blooming. Therefore, optimization of the blooming process in the in-line heat treatment materials was examined under laboratory experiments. As a result, it was found that, when the blooming is not executed, the grains tend to be fine in the in-line heat-treatment material, improving the toughness. Namely, it was found that the conventional general knowledge is not always correct.
In order to understand this unexpected result, a simulation test was further executed under laboratory experiments. In the process including the blooming process, a cast ingot was heated to 1250° C. and hot worked to form a block, which was then further heated to 1250° C. to perform hot rolling and water cooling, whereby the piercing process and the in-line heat treatment process were simulated.
In the process without including the blooming process, a block of the same size as the block formed by the above hot working was cut from the cast ingot by machining, and this block was heated to 1250° C. to perform hot rolling and water cooling, whereby the piercing process and the in-line heat treatment process were simulated.
As the results of the two simulation tests, the grain size not subjected to blooming was overwhelmingly fine and the toughness was improved.
However, the similar trial to the two simulation tests executed under actual equipments, not under laboratory experiments, resulted in the fact that the grains not subjected to blooming were not finer than expected.
Therefore, the inventors examined why the grain size not subjected to blooming differed extremely from the one subjected to blooming about the two simulation tests under laboratory experiments.
As a result, they found that most of added Ti precipitated as Ti carbonitride in the blooming simulation process under laboratory experiments, and that the number of precipitated grains reduced with the grain growing of the Ti carbonitride during heating and hot working in the blooming simulation process. The reduce of the number of precipitated grains deteriorated the capability of pinning the grain growth in the parent phase, which resulted in not suppressing the coarse-graining during the subsequent heating of the block for piercing simulation.
To the contrary, they found that, in the simulation test without the blooming process under laboratory experiments, the Ti carbonitride finely precipitated during the heating in the piercing process because no carbonitride precipitation generated within the ingot, and the Ti carbonitride pinned the grain growing in the parent phase, wherein remarkably fine grains were made.
The reason why the grains not subjected to blooming under the simulation tests under actual equipment were not finer than expected was found that the Ti carbonitride was already precipitated during the casting because the cooling rate during casting was not high enough to dissolve Ti in the solid state.
The Ti carbonitride that precipitated during casting is apt to coarsen with a reduced number of precipitated grains, since the precipitation is caused at a high temperature. Therefore, the capability of pinning the grains of the parent phase is reduced. On the other hand, if a sufficient quantity of dissolved Ti is ensured during the casting with minimized precipitation of Ti carbonitride, the Ti carbonitride is finely precipitated with an increased number of precipitated grains during the heating of the billet in the subsequent pipe making process because the precipitation is occurred at a low temperature. If the number of precipitated grains is large, the effect of pinning the crystal grains of the parent phase is increased to suppress the coarse-graining of the parent phase. Accordingly, it is extremely important to properly control the cooling rate during casting.
If the cooling rate after solidification is low, the Ti carbonitride precipitates in a high temperature range during cooling. However, this precipitate in an austenite range with relatively low dislocation causes few nucleation sites, which leads to a coarsely dispersed state. Once coarsely precipitated, the Ti carbonitride cannot be finely dispersed since it is hardly dissolved in a solid phase.
If the cooling rate after solidification is set to a rate causing no precipitation of the Ti carbonitride, the cast ingot has no Ti carbonitride, but Ti in a dissolved state. The Ti carbonitride precipitates at a relatively low temperature during the subsequent heating for hot working. Since, during heating, the Ti carbonitride precipitates at a low temperature in a bainite structure with high dislocation, the Ti carbonitride precipitates as being finely dispersed with many nucleation sites. It was also found that an excessively high heating rate makes fine precipitation difficult because of the precipitation in a high temperature range.
It is also effective for the sufficiently fine precipitation of Ti carbonitride to execute isothermal treatment in a proper temperature range during heating. The Ti carbonitride once finely precipitated is hardly coarsened, and even if blooming is executed, the effect of suppressing the coarse-graining can be exhibited. However, since a slight coarse-graining of Ti carbonitride is caused in the blooming, the dissolved Ti in solidification should preferably be present more than that in the case of executing without blooming.
Since precipitation strengthening by V or Nb makes easier to obtain high strength, the precipitation strengthening has been frequently applied to steel products, which require weldability in addition to high strength. However, it is better not to use the precipitation strengthening as much as possible, since it causes serious deterioration of toughness in a heavy wall in-line heat treatment material. Particularly, Nb seriously deteriorates the toughness of the in-line heat treatment material. Therefore, if Nb is included, it is necessary to strictly set an upper limit. With respect to V, it is also necessary to perform an alloy designing in order to ensure the strength based on the transformation strengthening and solid-solution strengthening by restricting the upper limit of the V content, although it is not as strict as in Nb.
Further, in the case of the heavy wall material, it is difficult to obtain a homogenous metal structure in quenching treatment during the first stage of the heat treatment, and the toughness tends to deteriorate. Since the cooling rate is reduced in the heavy wall material, it is difficult to obtain a homogenously transformed structure. Namely, although it is successively transformed to martensite or bainite during cooling, C is condensed to non-transformed austenite if the diffusion of C is possible to some degree with a low cooling rate, and this part is changed to martensite or bainite with a high C content or to retained austenite with a high C content after the final transformation. Accordingly, it is desirable to perform forced cooling to a temperature as low as possible, at a cooling rate as large as possible.
However, there are limits to increase the cooling rate in the case of heavy wall steel pipes. Therefore, examinations were made to develop a technique capable of forming a homogenous structure at a cooling rate that is attainable even in the heavy wall materials. Consequently, the inventors found that minimizing the content of the Carbon element to be condensed and also suppressing the content of Si can lead to reducing the condensation of C in the second phase.
Based on the above findings, basic ideas of the alloy design and the manufacturing process were clarified as follows to complete the present invention. In the following description, “%” represents “% by mass”, unless otherwise specified.
The C content is limited to not more than 0.08%. The upper limit of the Si content is set to not more than 0.25%, preferably to not more than 0.15%, and more preferably to not more than 0.10%. Ti content needs to be controlled in a narrow range of 0.004 to 0.010% suitable to precipitate as fine Ti carbonitride, without precipitation in the solidification, during the subsequent billet heating. Further, an addition of Nb is not performed in the case of an in-line heat treatment since it causes strength dispersion in addition to deterioration of the toughness, and the upper limit as an impurity is preferably set to not more than 0.005%. Since V also deteriorates the toughness, it is not added, or should be controlled to not more than 0.08% if included.
Other elements are adjusted from the viewpoint of the balance between high strength and satisfactory toughness. For P and S that adversely affect the toughness, the allowable upper limit values are set, respectively. Mn, Cr. Ni, Mo and Cu should be selectively adjusted according to the intended strength, considering the toughness and weldability. Al is added for deoxidation. It is also effective to selectively add at least one of Ca, Mg and REM to ensure the casting characteristic or improve the toughness. Further, the content of N needs to be controlled in a narrow range in order to precipitate stable Ti carbonitride.
For the manufacturing process, it is important to obtain a solidified ingot in which dissolved Ti is ensured while suppressing the precipitation of the Ti carbonitride. The present inventors found that the Ti carbonitride is not precipitated immediately after solidification if the contents of C, Ti, and N are set to the above ranges. However, since a coarse-grained Ti carbonitride is precipitated if the subsequent cooling rate is low, the cooling after solidification needs to be performed at a specified rate or more.
Regarding the casting method, a continuous casting to a billet with a circular cross section (hereinafter, refers to “a round billet”) is ideal, but a process of continuously casting to a square mold or casting thereto as an ingot, and then blooming to the round billet can be adapted. In this case, it is important to further strictly control the cooling rate after casting to ensure a sufficient quantity of dissolved Ti while suppressing precipitation of a coarse-grained TiN.
The round billet is reheated to a hot workable temperature and piercing, drawing and shaping rolling are performed thereto. If the dissolved Ti is sufficiently present, the Ti carbonitride is precipitated during reheating. Since the precipitation temperature is relatively low, remarkably fine Ti carbonitride is precipitated, compared with the precipitation during cooling after solidification. Since the number of grains of the finely precipitated Ti carbonitride is large, grain migration during heating or holding the billet can be suppressed to prevent the coarse-graining. A quick heating causes no fine precipitation at low temperature so that the effect of preventing the coarse-graining cannot be obtained. Therefore, a gentle heating or a holding in a middle stage is required to promote a precipitation of fine-grained Ti carbonitride.
To obtain a homogenous structure is necessary for ensuring the toughness in the heat treatment after pipe making. It is important therefore to use steel with an adjusted chemical composition and to sufficiently cool it at a forced cooling end temperature set as low as possible. These ideas result in improving the toughness because of preventing it from generating a transformation strengthened structure with partially concentrated C or retained austenite.
The present invention according to the above-mentioned basic ideas involves the following seamless steel pipes for line pipe (1) and (2) and the following methods of manufacturing a seamless steel pipe for line pipe (3) to (6).
(1) A heavy wall seamless steel pipe for line pipe with high strength and increased toughness, which has a chemical composition, by mass %, that consists of C: 0.03 to 0.08%, Si: not more than 0.25%, Mn: 0.3 to 2.5%, Al: 0.001 to 0.10%, Cr: 0.02 to 1.0%, Ni: 0.02 to 1.0%, Mo: 0.02 to 1.2%, Ti: 0.004 to 0.010%, N: 0.002 to 0.008%, and 0.0002 to 0.005%, in total, of at least one selected from Ca, Mg and REM, and the balance Fe and impurities, optionally including V: 0 to 0.08%, Nb: 0 to 0.05% or Cu: 0 to 1.0%, and that P and S among impurities are not more than 0.05% and not more than 0.005% respectively.
(2) A heavy wall seamless steel pipe for line pipe with high strength and increased toughness, which has 0.0003 to 0.01% of boron in addition to the chemical composition above.
(3) A method of manufacturing a heavy wall seamless steel pipe for line pipe with high strength and increased toughness characterized by comprising the following steps (a) to (e):
(a) Forming a billet with a round cross section by continuous casting of molten steel that has the chemical composition according to (1) or (2) above.
(b) Cooling the billet to the room temperature at not less than 6° C./min of an average cooling rate between 1400 and 1000° C.
(c) Heating the billet to the temperature between 1150 and 1280° C. at not more than 15° C./min of an average heating rate between 550 and 900° C., then piercing and rolling, to make a seamless pipe.
(d) Cooling forcedly the seamless pipe to a temperature of not higher than 100° C. at not less than 8° C./min of an average cooling rate between 800 and 500° C., immediately after pipe making, or after isothermal treating at the temperature between 850 and 1000° C. immediately in succession with pipe making, or after heating to the temperature between 850 and 1000° C. after once cooling in succession with pipe making.
(e) Tempering the seamless pipe at the temperature between 500 and 690° C.
(4) A method of manufacturing a heavy wall seamless steel pipe for line pipe with a high strength and increased toughness characterized by comprising the following steps (a) to (f):
(a) Forming a bloom or slab with a square cross section by continuous casting of molten steel that has the chemical composition according to (1) or (2) above.
(b) Cooling the bloom or slab to the room temperature at not less than 8° C./min of an average cooling rate between 1400 and 1000° C.
(c) Heating the bloom or slab to a temperature between 1150 and 1280° C. at not more than 15° C./min of an average heating rate between 550 and 900° C., and then cooling to the room temperature, to form a billet with a round cross section by forging and/or rolling.
(d) Heating the billet to the temperature between 1150 and 1280° C., then piercing and rolling, to make a seamless pipe.
(e) Cooling forcedly the seamless pipe to a temperature of not higher than 100° C. at not less than 8° C./min of an average cooling rate between 800 and 500° C., immediately after pipe making, or after isothermal treating at the temperature between 850 and 1000° C. immediately in succession with pipe making, or after heating to the temperature between 850 and 1000° C. after once cooling in succession with pipe making.
(f) Tempering the seamless pipe at the temperature between 500 and 690° C.
(5) A method of manufacturing a heavy wall seamless steel pipe for line pipe with a high strength and increased toughness characterized by comprising the following steps (a) to (e):
(a) Forming a billet with a round cross section by continuous casting of a molten steel that has the chemical composition according to (1) or (2) above.
(b) Cooling the billet to the room temperature at not less than 6° C./min of an average cooling rate between 1400 and 1000° C.
(c) Isothermal treating the billet during not less than 15 minutes at the temperature between 550 and 1000° C., and heating to the temperature between 1150 and 1280° C., then piercing and rolling, to make a seamless pipe.
(d) Cooling forcedly the seamless pipe to a temperature of not higher than 100° C. at not less than 8° C./min of an average heating rate between 800 and 500° C., immediately after pipe making, or after isothermal treating at the temperature between 850 and 1000° C. immediately in succession with pipe making, or after heating to the temperature between 850 and 1000° C. after once cooling in succession with pipe making.
(e) Tempering the seamless pipe at the temperature between 500 and 690° C.
(6) A method of manufacturing a heavy wall seamless steel pipe for line pipe with a high strength and increased toughness characterized by comprising the following steps (a) to (f):
(a) Forming a bloom or slab with a square cross section by continuous casting of a molten steel that has the chemical composition according to (1) or (2) above.
(b) Cooling the bloom or slab to the room temperature at not less than 8° C./min of an average cooling rate between 1400 and 1000° C.
(c) Isothermal treating the bloom or slab during not less than 15 minutes at the temperature between 550 and 1000° C., and heating to the temperature between 1150 and 1280° C., then forging and/or rolling, to form a billet with a round cross section, and then cooling to the room temperature.
(d) Heating the billet to the temperature of 1150 to 1280° C., then piercing and rolling, to make a seamless pipe.
(e) Cooling forcedly the seamless pipe to a temperature of not higher than 100° C. at not less than 8° C./min of an average cooling rate between 800 and 500° C., immediately after pipe making, or after isothermal treating at the temperature between 850 and 1000° C. immediately in succession with pipe making, or after heating to the temperature between 850 and 1000° C. after once cooling in succession with pipe making.
(f) Tempering the seamless pipe at the temperature between 500 and 690° C.
1. Chemical Composition of Steel Pipe of the Present Invention
The reason for limiting the chemical compositions of the steel pipe of the present invention will be described as follows. As described above, “%” showing the content (or concentration) of a chemical composition means “% by mass”.
C: 0.03 to 0.08%
C is an important element for ensuring the strength of steel. A content of not less than 0.03% is needed in order to improve the hardenability and strength in a heavy wall material. Since a content exceeding 0.08% causes deterioration of toughness, the C content is set to 0.03 to 0.08%.
Si: 0.25% or less
Si has an effect as a deoxidizer in steel production, but it is better to add as little as possible. Because it seriously deteriorates the toughness, particularly, of a heavy wall material. If the content of Si exceeds 0.25%, the toughness of the heavy wall material is remarkably deteriorated. Therefore, the content is set to 0.25% or less if it is added as the deoxidizer. A content of 0.15% or less enables further improvement in toughness. The content is most desirably controlled to less than 0.10%. Although it is difficult to extremely reduce Si as impurity from the point of the steel making process, extremely satisfactory toughness can be obtained if the content is limited to less than 0.05%.
Mn: 0.3 to 2.5%
Mn needs to be included in a relatively large quantity since it enhances the hardenability and therefore strengthens the center even in a heavy wall material and also enhances the toughness. A content of less than 0.3% cannot provide these effects, and a content exceeding 2.5% causes deterioration of the HIC resisting characteristic, therefore, the Mn content is set to 0.3 to 2.5%.
Al: 0.001 to 0.10%
Al is added as a deoxidizer in steel making. In order to obtain this effect, it needs to be added so as to have a content of not less than 0.001%. On the other hand, if the content of Al exceeds 0.10%, the inclusions are clustered, thereby deteriorating the toughness, and surface defects are frequently generated during the bevel face working of pipe ends. Therefore, the content of Al is set to 0.001 to 0.10%. From the point of preventing the surface defects, it is desirable to provide an upper limit, and the upper limit is preferably set to 0.03% and more preferably to 0.02%. Since a high deoxidation effect by the addition of Si cannot be expected in the steel pipe of the present invention, the lower limit of Al content is preferably set to 0.010% for sufficient deoxidation.
Cr: 0.02 to 1.0%
Cr is an element that improves the hardenability and strength of steel in a heavy wall material. The effect becomes remarkable when 0.02% or more of Cr is included. However, since an excessive content thereof causes some deterioration of toughness, the content is limited to 1.0% or less.
Ni: 0.02 to 1.0%
Ni is an element that improves the hardenability and strength of steel in a heavy wall material. The effect becomes remarkable when 0.02% or more of Ni is included. However, since Ni is an expensive element, and the effect is saturated if it is excessively included, the upper limit thereof is set to 1.0%.
Mo: 0.02 to 1.2%
Mo is an element that improves the strength of steel by transformation strengthening and solid-solution strengthening. The effect becomes remarkable when 0.02% or more of Mo is included. However, since an excessive addition thereof causes deterioration of the toughness, the upper limit is set to 1.2%.
Ti: 0.004 to 0.010%
The content of Ti needs to be controlled in a narrow range of 0.004% to 0.010% suitable to precipitate as fine Ti carbonitride, without precipitation in the solidification, during the subsequent billet heating. If the content is less than 0.004%, a sufficient number of precipitated grains of the Ti carbonitride cannot be ensured, and if it exceeds 0.010%, the Ti carbonitride is coarsely precipitated in the cooling after solidification. Therefore, a proper content of Ti is 0.004 to 0.010%.
N: 0.002 to 0.008%
N needs to be included in a content of 0.002% or more to ensure finely dispersed Ti carbonitride. Since a content exceeding 0.008% results in precipitation of coarse-grained Ti carbonitride in the solidification, the content needs to be controlled in a narrow range of 0.002 to 0.008%.
V: 0 to 0.08%
V is an element whose content is to be determined depending on the balance between strength and toughness. If sufficient strength can be ensured by other alloy elements, increased toughness can be obtained without the addition thereof. When it is added as a strength improving element, the content is preferably set to 0.02% or more. Since the toughness is seriously deteriorated if the content exceeds 0.08%, the upper limit of the content is set to 0.08% if added.
Nb: 0 to 0.05%
Nb is remarkably effective for suppressing the coarse-graining during heating for quenching in the case of an off-line heat treatment. In order to obtain this effect, Nb is desirably included in a content of 0.005% or more. However, if the content of Nb exceeds 0.05%, a coarse-grained carbonitride is precipitated to deteriorate the toughness. Therefore, the upper limit is set to 0.05%.
In the case of an in-line heat treatment, basically, it is better not to add Nb since the Nb carbonitride is inhomogeneously precipitated, which increases the strength dispersion as well as deterioration of the toughness. When the content exceeds 0.005%, the strength dispersion is remarkable and problematic in manufacturing. Therefore, when applying the in-line heat treatment, the allowable upper limit should be set to 0.005%.
Cu: 0 to 1.0%
Cu does not need to be added. However, it can be added, if improvement in the HIC resisting characteristic (hydrogen-induced cracking resisting characteristic) is intended, since it has an effect of improving the HIC resisting characteristic. The minimum content that improves the HIC characteristic is 0.02%. Since the effect is saturated even with a content that exceeds 1.0%, the content may be set to 0.02 to 1.0% if added.
Ca, Mg, and REM: 00002 to 0.005%, in total of at least one selected therefrom.
These elements are added for the purpose of improving the toughness and the corrosion resistance by shape controlling of inclusions and for the purpose of suppressing nozzle clogging in casting to improve the casting characteristic. In order to obtain such effects, a content of 0.0002% or more, in total of at least one selected therefrom is needed. If the content exceeds 0.005%, in total of at least one selected therefrom, not only the effects are saturated, but also the inclusions are easily clustered, thereby somewhat deteriorating the toughness and the HIC resisting characteristic. Therefore, when one of these elements are added, each content is set to 0.0002 to 0.005%, and when two or more selected therefrom are added, the total content is set to 0.0002 to 0.005%. REM means 17 elements that include lanthanoide elements, Y and Sc.
B: 0.0003 to 0.01%
B does not need to be added. However, since addition thereof leads to improvement in hardenability even if it is a trace, the addition is effective when further high strength is needed. In order to obtain this effect, a content of 0.0003% or more is desirable. However, since an excessive addition thereof causes deterioration of the toughness, the content of B is set to 0.01% or less if added.
The steel pipe for line pipe of the present invention contains the above chemical composition and the balance Fe and impurities. The upper limits of content of P and C among the impurities should be controlled as follows.
P: Not more than 0.05%
P is an impurity element that deteriorates the toughness, and the content is preferably as little as possible. Since a content exceeding 0.05% causes remarkable deterioration of the toughness, the allowable upper limit is set to 0.05%. The content of P is preferably 0.02% or less and, further preferably, 0.01% or less.
S: Not more than 0.005%
S is also an impurity element that deteriorates the toughness, and the content is preferably as little as possible. Since a content exceeding 0.005% causes remarkable deterioration of the toughness, the allowable upper limit is set to 0.005%. The content of S is preferably 0.003% or less and, further preferably, 0.001% or less.
2. Manufacturing Method
Suitable manufacturing conditions of the manufacturing method of the present invention will be described as follows.
(1) Casting and Cooling after Solidification
Steel is refined in a basic oxygen furnace or the like so as to have the above composition followed by casting and solidification to obtain a bloom. At this time, it is important to obtain a solidified ingot in which precipitation of the Ti carbonitride is suppressed. If the contents of C, Ti and N are restricted as described above, basically, the Ti carbonitride is not precipitated during solidification. However, if the subsequent cooling rate is low, a coarse-grained Ti carbonitride precipitates, therefore, the cooling must be performed at not less than a specified cooling rate.
A continuous casting to a round billet shape is ideal for the manufacturing process, however, a process for continuously casting to a square mold or casting thereto as an ingot and then blooming to the round billet can be performed. In this case, it is important to further strictly control the cooling rate after solidification to suppress the precipitation of a coarse-grained TiN.
An average cooling rate in a temperature range of 1400 to 1000° C., where the Ti carbonitride is apt to be generated after solidification, is required to be not less than 6° C./min in the case of casting to the round billet and to be not less than 8° C./min in the case of executing the blooming. The average cooling rate is more preferably set to be not less than 8° C./min in the case of casting to the round billet and to be not less than 10° C./min in the case of executing the blooming. In each case, no upper limit is provided since the larger average cooling rate is more desirable.
The cooling rate of the bloom varies depending on portions of the bloom. In the case of continuous casting to the circular mold, the cooling rate is controlled at a place distant from a center by the distance of ½ of the radius. In the case of continuously casting to a square mold, the cooling rate is controlled at a middle position between the center of gravity and the surface on a line passing the center of gravity of the square in parallel to the long sides thereof. The temperature can be measured by attaching a thermocouple, or instead, by numerical simulation corrected with the temperature history of the surface.
(2) Working of Billet or Ingot
The round billet is reheated to a hot workable temperature, and piercing, drawing and shaping rolling are performed thereto. The bloom or slab cast into a square cross section is reheated and then made into a round billet by forging or/and rolling, and piercing, drawing and shaping rolling are then performed thereto.
A reheating temperature is required to be 1150° C. or higher since hot deformation resistance is increased at a lower temperature than 1150° C., which increases flaws. The upper limit thereof is set to 1280° C., since a temperature exceeding 1280° C. leads to an excessive increase in the heating fuel raw unit, and a reduction in yield by increased scale loss, which uneconomically shortened the life of the heating furnace, and the like. Since a lower heating temperature makes finer grains and increases the toughness, a preferable heating temperature is 1200° C. or lower.
When the dissolved Ti is sufficiently present, the Ti carbonitride is precipitated during reheating, however, the precipitation occurs at a relatively low temperature, unlike the precipitation during the cooling after solidification. Therefore, the precipitated Ti carbonitride is much more fine-grained than the one during the cooling after solidification. An increased number of fine-grained Ti carbonitride are formed, and this suppresses the grain migration during heating the billet to prevent the coarse-graining. However, since rapid heating disables minute precipitation at a low temperature, the effect of preventing coarse-graining cannot be obtained. It is effective for the promotion of the minute precipitation at a low temperature during the reheating that the average heating rate should be 15° C./min or less at a temperature between 550 and 900° C. or that the isothermal treatment should be executed for 15 minutes or more at a temperature between 550 and 1000° C.
The piercing, drawing and shaping rolling can be executed in the manufacturing conditions for a general seamless steel pipe.
3. Heat Treatment after Pipe Making
In the heat treatment after pipe making, to obtain a homogeneous structure is needed in order to ensure toughness. The quenching treatment is based on the in-line heat treatment of executing quenching, without once cooling to room temperature, in succession with hot rolling. However, if the reheating and the quenching are performed after cooling, finer grains are made, improving the toughness. When the quenching is executed in succession with isothermal treatment in an isothermal furnace after the end of hot work, a steel pipe with minimized strength dispersion can be obtained.
High strength and high toughness can be obtained more easily in a heavy wall material if the cooling rate in the quenching is set higher. As the cooling rate gets closer to a theoretically limited cooling rate, higher strength and higher toughness can be obtained. The necessary average cooling rate is 8° C./sec or more at a temperature between 800 and 500° C., more preferably, 10° C./sec or more, and most preferably not less than 15° C./sec.
The cooling end temperature is also important for ensuring excellent toughness, in addition to the cooling rate. It is important to use a steel with an adjusted chemical composition and forcedly cool it to an end temperature of not higher than 100° C. The forced cooling is continuously performed, preferably to 80° C. or lower, more preferably to 50° C. or lower, and most preferably to 30° C. or lower. According to this, the generation of transformation strengthened structure in which C is partially concentrated or retained austenite can be prevented, therefore, the toughness is significantly improved.
After the quenching, tempering is performed at a temperature between 500 and 700° C. The tempering is performed in order to adjust the strength and improving the toughness. The holding time at the tempering temperature may be properly determined according the wall thickness or the like of the steel pipe, and is generally set to about 10 to 120 minutes.
Steels that have the chemical compositions shown in Table 1 were molten in a converter. Two methods for manufacturing a round billet were adopted: one is a method for casting to a continuous mold with a round cross section, and the other is a manufacturing method for casting to a square mold and then blooming. The manufacturing conditions in the casting to the round continuous casting mold are shown in Tables 2 and 3. The solidification process is represented as “RCC”. The process for casting to the square mold is represented as “BLCC”, and the manufacturing conditions thereof are shown in Tables 4 and 5.
Round billets were heated in the pipe making heating conditions shown in Tables 2 to 5, and hollow pipes were produced by use of a feed roll piercing machine. The hollow pipes were finish-rolled by use of a mandrel mill and a sizer, whereby steel pipes having wall thickness of 30 mm to 50 mm were obtained. Thereafter, these pipes were cooled in the quenching conditions described in Tables 2 to 5. Namely, after pipe making, any one of the following three processes were adopted: the first one is cooling immediately; the second one is charging immediately to a reheating furnace for isothermal treatment and then quenching; and the last one is cooling once to room temperature and reheating and then cooling again. Thereafter, tempering was executed in the conditions described in Tables 2 to 5 to obtain the finished products.
A JIS (Japan Industrial Standard) No. 12 tensile test piece for tensile test was prepared from each of the resulting steel pipes to measure tensile strength (TS) and yield strength (YS). The tensile test was carried out according to JIS Z 2241. As an impact test piece, a V notch test piece of 10 mm×10 mm, 2 mm was prepared from the longitudinal direction of the heavy wall center according to No. 4 test piece of JIS Z 2202, and subjected to the test.
In Test No. 1 of Table 2, two examples of branch numbers of 1 and 2 are described. Steel A, the invention, is used in 1-1 and 1-2, and the manufacturing condition of 1-1 is within the range restricted by the present invention, where increased toughness is obtained. On the other hand, the manufacturing condition of 1-2 is deviated from the manufacturing process defined by the present invention with an excessively high heating rate for pipe making, where increased toughness cannot be obtained. Each of Test Nos. 2 to 24 also has branch numbers 1 and 2, and the same steel grade is used in the same test number. The manufacturing condition of each branch number 1 is within the range restricted by the present invention, where increased toughness can be obtained. On the other hand, the manufacturing condition of each branch number 2 is deviated from the manufacturing process defined by the present invention, where increased toughness cannot be obtained.
In Tables 4 and 5, the same steel grade is also used in one test number, and each branch number 1 corresponds to the manufacturing process within the range restricted by the present invention, where increased toughness is obtained. On the other hand, since each branch number 2 is deviated from the manufacturing process defined by the present invention, where increased toughness is not obtained.
Test Nos. 25 to 30 are examples of comparative steels that are deviated from the alloy composition range restricted by the present invention. Each of the steels is insufficient in toughness, and has insufficient performances as a line pipe requiring increased thickness and high toughness.
(*) Note: The symbol I shows the invention and the symbol C shows the comparative.
(*) Note: The symbol I shows the invention and the symbol C shows the comparative.
(*) Note: The symbol I shows the invention and the symbol C shows the comparative.
(*) Note: The symbol I shows the invention and the symbol C shows the comparative.
According to the present invention, by restricting the chemical composition of a seamless steel pipe and the manufacturing method thereof, a seamless steel pipe, even a heavy wall steel pipe, for line pipe excellent in toughness while having high strength such as yield strength of X70 class (yield strength of not less than 482 MPa), X80 class (yield strength of not less than 551 MPa), X90 class (yield strength of not less than 620 MPa), X100 class (yield strength of not less than 689 MPa), and X120 class (yield strength of not less than 827 MPa) can be manufactured. The seamless steel pipe of the present invention is a steel pipe that can be laid in a severer circumstance of the deep sea, particularly, for the use of a submarine flow line. The present invention is thus greatly contributable to stable supply of energies.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2005-095240 | Mar 2005 | JP | national |
| Number | Date | Country | |
|---|---|---|---|
| Parent | PCT/JP2006/304613 | Mar 2006 | US |
| Child | 11895131 | Aug 2007 | US |
