HIGH-STRENGTH SEAMLESS STEEL TUBE AND METHOD FOR MANUFACTURING IT

Abstract

The present invention provides a high-strength seamless steel tube and a method for manufacturing the same. A high-strength seamless steel tube of the present invention has an inflow of 758 MPa or more, and a KILOMETER value of 23.0 MPavm or more according to a sulfide stress corrosion cracking resistance evaluation index.

Description

HIGH-STRENGTH SEAMLESS STEEL TUBE AND METHOD FOR MANUFACTURING IT FIELD OF INVENTION The present invention relates to a high-strength seamless steel pipe for oil and gas wells, specifically, a high-strength seamless steel pipe that has excellent resistance to sulfide stress corrosion cracking (SSC resistance) in acidic environments containing hydrogen sulfide. The present invention also relates to a method for manufacturing such a high-strength seamless pipe. BACKGROUND OF THE INVENTION The increase in crude oil prices and the anticipated scarcity of oil resources in the near future have spurred the active development of oil and gas fields that were previously inconceivable, such as deep-water oil fields and oil and gas fields in severely corrosive environments containing hydrogen sulfide, or acidic environments as they are also called. Steel pipes for oilfield tubular products used in these environments are required to be made of materials with high strength and superior corrosion resistance (acid resistance). In response to those demands, for example, PTL 1 describes a tubular products steel for oilfields having improved sulfide stress corrosion cracking resistance, specifically, a low-alloy steel comprising, in wt.%, C: 0.2 to 0.35%, Cr: 0.2 to 0.7%, Mo: 0.1 to 0.5%, and V: 0.1 to 0.3%, and specifying a total amount of precipitating carbides, and the fraction of MC-type carbides therein. PTL 2 describes a steel material for oilfield tubular products that has improved sulfide stress corrosion cracking resistance. The steel material described in this prior art document comprises, in % by mass, C: 0.15 to 0.30%, Si: 0.05 to 1.0%, Mn: 0.10 to 1.0%, P: 0.025% or less, S: 0.005% or less, Cr: 0.1 to 1.5%, Mo: 0.1 to 1.0%, Al: 0.003 to 0.08%, N: 0.008% or less, B: 0.0005 to 0.010%, and Ca+O (oxygen): 0.008% or less, and one or two or more selected from Ti: 0.005 to 0.05%, Nb: 0.05% or less, Zr: 0.05% or less, and V: 0.30% or less less. Regarding the properties of inclusions in steel, the steel specifies the maximum length of non-continuous metallic inclusions, and the number of particles with a particle diameter of 20 pm or more. PTL 3 describes a tubular steel for oilfield products that has improved resistance to sulfide stress corrosion cracking. The steel described in this related prior art document comprises, in mass percent, C: 0.15 to 0.35%, Si: 0.1 to 1.5%, Mn: 0.1 to 2.5%, P: 0.025% or less, S: 0.004% or less, sol.Al: 0.001 to 0.1%, and Ca: 0.0005 to 0.005%, and specifies the composition of non-metallic inclusions based on Ca, the compound Ca-Al oxide, and the HRC hardness of the steel. PTL 4 describes a low-alloy steel for oilfield tubular products that has improved resistance to sulfide stress corrosion cracking and a yield strength of 861 MPa or higher. The low-alloy steel described herein, in relation to the prior art, comprises, in mass percent, C: 0.2 to 0.35%, Si: 0.05 to 0.5%, Mn: 0.05 to 1.0%, P: 0.025% or less, S: 0.01% or less, Al: 0.005 to 0.10%, Cr: 0.1 to 1.0%, Mo: 0.5 to 1.0%, Ti: 0.002 to nrnonn / zznz / B / YiAi 0.05%, V: from 0.05 to 0.3%, B: from 0.0001 to 0.005%, N: 0.01% or less, and O: 0.01% or less, and establishes a default value for a formula containing the entire width of half the maximum of the

[211] plane of the steel, and a hydrogen diffusion coefficient. The sulfide stress corrosion cracking resistance of the steels described in PTL 1 through PTL 3 is a measure of the presence or absence of SSC after a round bar tensile test specimen is immersed in a test bath under a constant tensile load for 720 hours in compliance with NACE (National Association for Corrosion Engineering) Method A TM0177. The sulfide stress corrosion cracking resistance of the steel described in PTL 4 is a measure of whether the Kissc stress intensity factor value obtained in a hydrogen sulfide corrosive environment after the DCB (Double Cantilever Beam) test performed in compliance with NACE TM0177 Method D is equal to or greater than a specified value. List of Appointments Patent Literature PTL 1: JP-A-2000-178682 PTL 2: JP-A-2001 -172739 PTL 3: JP-A-2002-60893 PTL 4: JP-A-2005-350754 BRIEF DESCRIPTION OF THE INVENTION Technical Problem The 2016 revisions to NACE TM0177 introduced the Kilolimit value, a new index of resistance to sulfide stress corrosion cracking. Figure 1 is a diagram illustrating the method for determining a Kilolimit value. To determine a Kilolimit value, the applied stress intensity factor Kiapiicada at the tip of a notch on a test specimen before the start of a DCB test is plotted against the Kissc value obtained from a DCB test performed multiple times under different test conditions, as shown in Figure 1. A Kilolimit value can then be determined from the intersection of the linear regression line of the Kissc values ​​and the line where Kissc and Kiapiicada are one-to-one. In Figure 1, the vertical and horizontal axes represent Kissc and Kiapiicada, respectively.From PTL 1 to PTL 4, nothing is described about the specific measurements to improve the Kilolimit value to ensure resistance to sulfide stress corrosion cracking using the Kilolimit value. The present invention was developed in response to the problems discussed above, and one objective of the present invention is to provide a high-strength seamless steel pipe having high strength equivalent to at least API C110 standards, and having excellent sulfide stress corrosion cracking resistance (SSC resistance), specifically, a high and stable Kilolimit value, in acidic environments containing hydrogen sulfide. The present invention is also intended to provide a method for manufacturing such a high-strength seamless steel pipe. Solution to the Problem The inventors of the present invention conducted intensive studies to find a solution to the above problems. First, three types of steel pipe materials (steel Nos. A to C) were prepared that had compositions meeting API standards, as shown in Table 1. These steel pipe materials were used to produce test steel pipes (steel pipes without ncnonn / zznz / B / YiAi Table 1 Steel No. Composition (% by mass) C Si Mn PS Cr Mo Al Cu Nb VB Ti O N Ca A 0.24 0.03 0.59 0.006 0.0005 0.96 0.66 0.068 0.06 0.033 0.038 0.0019 0.001 0.0009 0.0025 - B 0.26 0.04 0.56 0.005 0.0006 1.04 0.73 0.066 0.05 0.031 0.032 0.0022 0.002 0.0010 0.0028 0.0011 C 0.25 0.19 0.51 0.008 0.0007 0.92 0.80 0.067 0.07 0.026 0.027 0.0027 0.003 0.0012 0.0037 - ncnonn / zznz / B / YiAi Specifically, several types of blocks were used for the hot rolling experiment, representing the three types of steel pipe materials used to form the test tubes. The block was tested in a plate rolling and direct quenching experiment that simulates the hot forming and subsequent direct quenching of a seamless steel pipe, using a small-scale hot rolling mill, a cooling device, and a heating furnace. After adjusting the creep of the rolled plates to a value equivalent to API C110 by quenching and tempering, a DCB test specimen was taken from the material and tested using a DCB test. The test was performed under the same conditions described above. The Kilolimit value obtained in the DCB test was examined for any relationship with various rolling conditions.As a result, it was found that the Kilílimit value improves particularly with the decrease in the starting temperatures of the intermediate heating carried out after the piercing and elongation rolling and before the sizing rolling of the seamless steel tube. The inventors herein conducted further research. Figure 5 depicts seamless steel tube manufacturing processes. As shown in Figure 5, the inventors herein devised a modification to a traditional seamless steel tube manufacturing process by adding intermediate cooling before the intermediate heating performed after the punching and lengthening roll and before the sizing roll. It was found that the cooling interruption temperature (specifically, the recovery temperature after intermediate cooling; described below) and the time before the subsequent intermediate heating begins are important in the intermediate cooling process. To investigate this, the inventors of the present work carried out a plate rolling and direct quenching experiment simulating the hot forming and subsequent direct quenching of a seamless steel tube, and performed intercooling during plate rolling. In the experiment, the recovery temperature after intercooling and the time before the start of intercooling were varied. Separately, a sample prepared by reheat quenching and tempering of the rolled material was subjected to a DCB test, and the Kilolimit value obtained in the test was used to find the optimum combination of recovery temperature after intercooling and time before the start of intercooling. Figure 7 is a diagram representing the Kilolimit values ​​classified on the graph of waiting time tW before the start of intermediate heating (seconds) plotted against (Tr-Ms)(°C), a value obtained by subtracting the martensitic transformation temperature Ms (°C) of the sample from the recovery temperature Tr (°C) after intermediate cooling. In Figure 7, the open circle represents the experimental conditions that produced a target Kilolimit value of 23.0 MPa or higher, and the cross represents the experimental conditions under which the Kilolimit value was below the target value of 23.0 MPa. Kilolimit was found not to meet the target value when the recovery temperature Tr (°C) after intermediate cooling exceeds (Ms+150°C), despite the waiting time tW before the start of intermediate heating.One possible explanation for this observation is that, even with intermediate cooling, the transformation (likely the bainite transformation) does not occur after cooling and before the start of intermediate heating when the cooling interruption temperature (specifically, the recovery temperature after intermediate cooling, described above) exceeds (Ms+150°C). It was also found that Kilílimit can more easily satisfy the target value as the recovery temperature Tr after intermediate cooling decreases, even when the waiting time tW before the start of intermediate heating is short, as shown in Figure 7.Presumably, with intermediate cooling, the bainite transformation begins when the recovery temperature Tr after intermediate cooling is (Ms+150°C) or lower, and proceeds during the waiting time before the start of intermediate heating, allowing the reverse transformation to occur during subsequent intermediate heating. The resulting grain refining appears to be the reason for the improved Kilolimit value. The present invention was completed on the basis of those discoveries, and the essence of the present invention is as follows. [1] A high-strength seamless steel pipe having a yield strength of 758 MPa or more, and a Kilolimit value of 23.0 MPa / mo plus as an index of evaluation of resistance to sulfide stress corrosion cracking. Here, Kilílimit is the determined value of the intersection between (i) a linear regression line created by a stress intensity factor Kissc obtained in a DCB (Double Cantilever Beam) test performed multiple times under different test conditions, and a stress intensity factor applied Kiapiicada at the tip of a notch in a test specimen prior to the start of the DCB test, and (ii) a straight line on which Kissc and Kiapiicada are one to one. [2] High-strength seamless steel pipe according to [1], having a steel microstructure with a pre-austenitic grain size of 10.5 or more in terms of a grain size number in compliance with ASTM E112. [3] High-strength seamless steel tube according to [1] or [2], having a composition that includes, in % by mass, C: 0.23 to 0.27%, Si: 0.35% or less, Mn: 0.45 to 0.70%, P: 0.010% or less, S: 0.0010% or less, Cr: 0.80 to 1.20%, Mo: 0.50 to 0.90%, Al: 0.080% or less, Cu: 0.09% or less, Nb: 0.050% or less, V: 0.050% or less, B: 0.0015 to 0.0030%, Ti: 0.005% or less, O: 0.0020% or less, and N: 0.0050% or less, and in of which the remainder is Fe and accidental impurities. [4] High-strength seamless steel tube according to [3], wherein the composition further includes, in % by mass, Ca: 0.0020% or less. [5] A method for manufacturing the high-strength seamless steel tube of any of [1] to [4], the method includes: a heating step of a steel tube material to a heating temperature ncnonn / zznz / B / YiAi in a temperature region of 1,200 to 1,300°C; a first hot rolling step of hot rolling of the hot steel tube material by punching and lengthening the steel tube material with a final rolling temperature of 800°C or more; an intermediate cooling step of cooling a raw steel tube after the first hot rolling step, the raw steel tube being cooled from a cooling start temperature of 700°C or more under conditions where the average cooling rate is 40°C / s or more, and the recovery temperature Tr of the raw steel tube at a tube surface is (Ms+150°C) or less, where Ms is a martensitic transformation start temperature; an intermediate heating step of heating the raw steel tube after the intermediate cooling step, the raw steel tube being heated to a surface temperature of 800 to 1,000°C after a holding time tW of 300 seconds or less being loaded into a reheating furnace; a second hot rolling step of subjecting the raw steel tube after the intermediate heating step to sizing hot rolling at a temperature equal to or greater than (Ar3+100°C), where Ar3 is a ferrite transformation start temperature, and ending the hot rolling at a temperature of (Ar3+50°C) or higher; a direct quenching step of directly quenching the raw steel tube continuously from the second hot rolling step, the raw steel tube being quenched to a temperature equal to or greater than (Ar3+10°C) under conditions where the average cooling rate is 40°C / s more, and the cooling interruption temperature is 200°C or less; and a heat treatment step of subjecting the raw steel tube after the direct quenching step to at least one heat treatment run quenching the raw steel tube after reheating to a temperature of 850 to 930°C, and continuously quenching the raw steel tube by heating from 650 to 730°C, satisfying the recovery temperature Tr and the holding time tW in the intermediate heating step according to the following formula (1): (Tr-Ms) < 10 + 0.0024 x (tW)2... (1). As used herein, “high strength” means strength equivalent to at least the C110 grade of the API standards, specifically, strength with a yield strength of 758 MPa or more (110 ksi or more). A high-strength seamless steel tube of the present invention has excellent sulfide stress corrosion cracking resistance (SSC resistance). Herein, “excellent sulfide stress corrosion cracking resistance” means that it has a Kilolimit value of 23.0 MPa / m or more as calculated using the method of Figure 1, using the Kissc (MPa / m) obtained by varying the wedge thickness in a DCB test performed according to NACE TM0177 Method D with a test bath using an aqueous solution at 24°C of 5% by mass of NaCl and 0.5% by mass of CH3COOH saturated with 1 atm (0.1 MPa) of hydrogen sulfide gas. Advantageous Effects of the Invention The present invention can provide a high-strength seamless steel pipe having high strength equivalent to at least API C110 grade, and excellent resistance to sulfide stress corrosion cracking (SSC resistance), specifically a high Kilolimit value, in acidic environments containing hydrogen sulfide. The present invention can also provide a method for manufacturing such a high-strength seamless steel pipe. BRIEF DESCRIPTION OF THE FIGURES Figure 1 is a diagram that represents a method for deriving the Kilímite value. Figure 2 is a diagram representing the shape and dimensions of a DCB test specimen. Figure 3 is a diagram representing the shape and dimensions of a wedge used in the DCB test. Figure 4 is a diagram that represents the relationship between the yield strength (YS) and Kilolimit value of seamless steel pipe for different seamless steel pipe manufacturing processes. Figure 5 is a diagram comparing a traditional seamless steel pipe manufacturing process and a seamless steel pipe manufacturing process of the present invention. Figure 6 is a diagram that represents the time-dependent temperature changes on the outer surface, center of wall thickness, and inner surface of a raw steel tube according to the heat transfer calculations of a water-cooled raw steel tube for seamless steel tubes. Figure 7 is a diagram representing the measurement result of Kilolimit values ​​obtained for experimental materials simulating seamless steel tubes and plotted on a graph of recovery temperature after intermediate cooling with water and waiting time before the start of intermediate heating after recovery. DETAILED DESCRIPTION OF THE INVENTION The following specifically describes the present invention. It should be noted that the present invention is not limited to the following embodiments. First, a high-strength seamless steel tube of the present invention is described. A high-strength seamless steel tube of the present invention has a yield strength of 758 MPa or more, and a Kilolimit value of 23.0 MPa or more as an index for evaluating resistance to sulfide stress corrosion cracking. Herein, Kilolimit is a value determined from the intersection between (i) a linear regression line created by the stress intensity factor Kissc obtained in a DCB (Double Cantilever Beam) test performed multiple times under different test conditions, and the applied stress intensity factor Kiapiicada at the tip of a notch in a test specimen before the start of the DCB test, and (ii) a straight line on which Kissc and Kiapiicada are equal. As mentioned above, a high-strength seamless steel pipe of the present invention has a high strength equivalent to at least API C110 (a yield strength of 758 MPa or greater) and excellent sulfide stress corrosion cracking (SSC) resistance in acidic environments containing hydrogen sulfide. Here, the yield strength is 758 MPa or greater, and the Kilolimit value is 23.0 MPa or greater, following the discussion provided above and the detailed descriptions of the reasons why those specific values ​​are omitted. The yield strength is preferably less than 862 MPa. The target Kilolimit value is set at 23.0 MPa or greater based on the maximum anticipated notch defect and the load application conditions of the tubular products for oilfields. The target value for Kilolimit is preferably 24.0 MPa / mo more, more preferably 25.0 MPa / mo more. Preferably, a high-strength seamless steel tube of the present invention has a steel microstructure with a pre-austenitic grain size of 10.5 or more in terms of a grain size number in compliance with ASTM E112 (hereafter referred to as “pre-austenitic grain size”). A pre-austenitic grain size of less than 10.5 leads to insufficient grain refinement, and the Kilolimit may fail to meet its target value. For this reason, the pre-austenitic grain size is preferably 10.5 or larger. More preferably, 11.0 or larger, and even more preferably 12.0 or larger. From the perspective of grain refinement limits in actual production, the pre-austenitic grain size is preferably 17.0 or smaller. The size of previous austenitic grains can be measured using the method described in the Examples of the present invention below. The following describes the preferred composition ranges for the high-strength seamless steel tube of the present invention, along with the reasons for the preferred ranges. In the following, “%” is the mass percentage (% by mass), unless specifically stated otherwise. C: from 0.23 to 0.27% Carbon acts to increase the strength of steel and is preferably contained in an amount of 0.23% or more to achieve high strength with a yield strength of 758 MPa or more. A carbon content of more than 0.27% considerably hardens the steel and can lead to a deterioration of the ultimate tensile strength (Kilometer value). For this reason, the carbon content is preferably 0.23 to 0.27%. The carbon content is more preferably 0.24% or more. The carbon content is more preferably 0.26% or less. Yes: 0.35% or Less Si is an element that acts as a deoxidizing agent, suppressing abrupt softening during tempering and increasing the steel's strength by forming a solid solution within the steel. It is preferably contained in an amount of 0.01% or more to achieve these effects. A Si content exceeding 0.35% can lead to the formation of coarse oxide inclusions and a deterioration of the Kilolimit value. For this reason, the Si content is preferably 0.35% or less. More preferably, the Si content is 0.01% or more, and even more preferably 0.02% or more. The Si content is also more preferably 0.20% or less, and even more preferably 0.04% or less. Mn: from 0.45 to 0.70% Manganese (Mn) is an element that increases the strength of steel by improving its hardenability. It also acts to fix sulfur by forming MnS with S, preventing sulfur-induced embrittlement at grain boundaries. In the present invention, Mn is preferably contained in an amount of 0.45% or more. An Mn content of more than 0.70% can considerably harden the steel as a result of the improved hardenability and may lead to a deterioration of the Kilolimit value. For this reason, the Mn content is preferably from 0.45% to 0.70%. The Mn content is more preferably 0.50% or more, and even more preferably [unclear text - likely a reference to a specific type of Mn]. 0.55% or more. The Mn content is more preferably 0.65% or less, even more preferably 0.60% or less. P: 0.010% or less P can segregate at grain boundaries or other parts of the steel in a solid solution state, causing defects such as grain boundary embrittlement cracking. In the present invention, P is preferably contained in as small an amount as possible, and a P content as low as 0.010% is acceptable. For this reason, the P content is preferably 0.010% or less. The P content is more preferably 0.008% or less, and even more preferably 0.006% or less. S: 0.0010% or less Sulfur exists almost entirely as sulfide inclusions in steel and can decrease ductility, toughness, and corrosion resistance, such as resistance to sulfide stress corrosion cracking. Sulfur can also exist partially in solid solution. In this case, the sulfur segregates at grain boundaries and other parts of the steel and can cause defects such as grain boundary embrittlement cracking. For this reason, in the present invention, sulfur is preferably contained in as small an amount as possible. However, excessive reduction of the S content leads to high refining costs. Therefore, in the present invention, the S content is preferably 0.0010% or less, a range in which the adverse effect of this element is tolerable. The S content is more preferably 0.0008% or less, and even more preferably 0.0006% or less. Cr: 0.80 to 1.20% Cr is an element that contributes to increasing the strength of steel by increasing its hardenability and improving corrosion resistance. Cr also forms carbides such as M3C, M7C3, and M23C6 by bonding with carbon during quenching, and these carbides, M3C carbide in particular, improve resistance to quench softening. In this way, Cr reduces strength fluctuations due to quenching and contributes to improved creep. Cr is preferably contained in an amount of 0.80% or more to achieve a creep of 758 MPa or more. A Cr content of more than 1.20% is economically disadvantageous because the effect becomes saturated. For this reason, the Cr content is preferably between 0.80% and 1.20%. The Cr content is more preferably 0.90% or more, and even more preferably 0.95% or more. The Cr content is more preferably 1.10% or less, even more preferably 1.05% or less. Mo: from 0.50 to 0.90% Molybdenum (Mo) is an element that contributes to increasing the strength of steel by increasing its hardenability and improving its corrosion resistance. Molybdenum, particularly in the form of M0₂C carbides formed through secondary precipitation after quenching, improves resistance to quench softening. In this way, molybdenum reduces strength variations due to quenching and contributes to improved creep. Mo is preferably contained in an amount of 0.50% or more to obtain these effects. A Mo content of more than 0.90% is economically disadvantageous because the effect becomes saturated. For this reason, the Mo content is preferably between 0.50% and 0.90%. The Mo content is more preferably 0.60% or more, and even more preferably 0.65% or more. The Mo content is most preferably 0.80% or less, and even more preferably 0.75% or less. ncnonn / zznz / B / YiAi Al: 0.080% or less Aluminum acts as a deoxidizing agent and contributes to reducing nitrogen in the solid solution by forming alkali metal ions (AmN) with nitrogen. To achieve this effect, aluminum is preferably present at a concentration of 0.040% or more. An aluminum content greater than 0.080% can increase oxide inclusions and may lead to a deterioration of the Kilolimit value. For this reason, the aluminum content is preferably 0.080% or less. More preferably, the aluminum content is 0.050% or more. Even more preferably, the aluminum content is 0.070% or less. Cu: 0.09% or less Copper (Cu) is an element that enhances corrosion resistance. When added in trace amounts, Cu forms dense corrosion products and suppresses the generation and growth of pitting, which can become the starting points for sulfide stress cracking (SSC). In this way, Cu greatly improves resistance to sulfide stress cracking. For this reason, in the present invention, Cu is preferably contained in an amount of 0.02% or more. A Cu content of more than 0.09% can lead to a decrease in hot workability during the seamless steel pipe manufacturing process. Therefore, the Cu content is preferably 0.09% or less. The Cu content is more preferably 0.03% or more, and even more preferably 0.04% or more. The Cu content is most preferably 0.07% or less, and even more preferably 0.06% or less. Nb: 0.050% or less Nitrogen (Nb) is an element that contributes to the refinement of γ grains by retarding recrystallization in the austenite (γ) temperature region, and it is very effective in refining substructures (e.g., bundles, blocks, and laths). Nb also strengthens steel by forming carbides. To achieve these effects, Nb is preferably present at a concentration of 0.020% or more. An Nb content greater than 0.050% promotes the formation of coarse precipitates (NbN) and can lead to a deterioration of the Kilolimit value. For this reason, the Nb content is preferably 0.050% or less. More preferably, the Nb content is 0.025% or more, and even more preferably 0.030% or more. The Nb content is also preferably 0.045% or less, and even more preferably 0.040% or less.Here, the “package” is defined as a region formed by aggregates of strips that have parallel faces with the same usual plane, while the “block” is formed by aggregates of parallel strips of the same orientation. V: 0.050% or less Zinc (V) is an element that forms carbides or nitrides and contributes to the strengthening of steel. To achieve these effects, it is preferably contained in an amount of 0.020% or more. A V content greater than 0.050% results in the thickening of the V carbides, which become points of initiation for sulfide stress corrosion cracking, and decreases the Kilolimit value rather than increasing it. For this reason, the V content is preferably 0.050% or less. The V content is more preferably 0.025% or more, and even more preferably 0.030% or more. The V content is most preferably 0.045% or less, and even more preferably 0.040% or less. B: from 0.0015 to 0.0030% Boron (B) is an element that contributes to improved hardening capacity when present in trace amounts. In the present invention, B is preferably contained in an amount of ncnonn / zznz / B / YiAi of 0.0015% or more. A boron content of more than 0.0030% is economically disadvantageous because the effect becomes saturated, or the desired effect cannot be expected as a result of the formation of iron boride (Fe-B). For this reason, the B content is preferably from 0.0015% to 0.0030%. The B content is more preferably 0.0016% or more, and even more preferably 0.0018% or more. The B content is most preferably 0.0027% or less, and even more preferably 0.0023% or less. Ti: 0.005% or less Ti forms nitrides and can cause deterioration of the Kilolimit value as a result of coarse titanium nitride becoming a point of origin for SSC. For this reason, the Ti content is preferably 0.005% or less. The Ti content is more preferably 0.003% or less. The Ti content is even more preferably 0.002% or less. O (Oxygen): 0.0020% or less In steel, oxygen (O₂) exists as accidental impurities in the form of oxides of elements such as aluminum (Al) and silicon (Si). Oxygen can cause deterioration of the Kilolimit value when coarse oxides are present in large quantities. For this reason, the oxygen content is preferably 0.0020% or less, a range in which the adverse effect of this element is tolerable. More preferably, the oxygen content is 0.0015% or less, and even more preferably, 0.0010% or less. N: 0.0050% or less Nitrogen (N) represents incidental impurities in the steel and forms MN-type precipitates by bonding with nitride-forming elements such as Al, Nb, and Ti. Excess nitrogen from the formation of these nitrides bonds with boron to form BN precipitates. Because this eliminates the hardenability-enhancing effect of boron addition, the amount of excess nitrogen should preferably be reduced as much as possible, ideally to 0.0050% or less. The N content is more preferably 0.0040% or less, and even more preferably 0.0030% or less. In the composition of the above components, the remainder is Fe and accidental impurities. Preferably, the high-strength seamless steel tube of the present invention contains the above components as its basic composition. The desired properties of the present invention can be obtained with the above preferred elements. Optionally, it may contain 0.0020% or less calcium for further improvement of strength and SSC resistance. Ca: 0.0020% or less Calcium (Ca) is effective in preventing nozzle clogging during continuous casting and is contained in a desirable amount of 0.0005% or more to achieve the desired effect. As an alternative to manganese (Mn), Ca fixes sulfur by forming CaS with S, preventing grain boundary embrittlement caused by sulfur. Unlike MnS, which is ductile, calcium disperses finely in the steel without elongation during hot rolling and improves resistance to sulfide stress corrosion cracking. However, Ca forms non-metallic oxide inclusions by combining with aluminum (Al), and when contained in an amount particularly greater than 0.0020%, calcium forms these inclusions in large quantities, causing deterioration of the Kilolimit value. For this reason, when Ca is included, it is preferably contained in an amount of 0.0020% or less. The Ca content is more preferably 0.0007% or more, and even more preferably 0.0007%.0.0009% or more. The Ca content is more preferably 0.0015% or less, even more preferably 0.0012% or less. The following describes a method for manufacturing high-strength seamless steel tubing of an embodiment of the present invention. A method for manufacturing high-strength seamless steel tubing of the present invention includes: a heating step of a steel pipe material of the above specific composition to a heating temperature in a temperature region of 1,200 to 1,300°C; a first hot rolling step of hot rolling of the hot steel tube material by punching and lengthening the steel tube material with a final rolling temperature of 800°C or more; an intermediate cooling step of cooling a raw steel tube after the first hot rolling step, the raw steel tube being cooled from a cooling start temperature of 700°C or more under conditions where the average cooling rate is 40°C / s or more, and the recovery temperature Tr of the raw steel tube at a tube surface is (Ms+150°C) or less, where Ms is the martensitic transformation start temperature calculated from formula (A) below; an intermediate heating step of heating the raw steel tube after the intermediate cooling step, the raw steel tube being heated to a surface temperature of 800 to 1,000°C after a holding time tW of 300 seconds or less being loaded into a reheating furnace; a second hot rolling step of subjecting the raw steel tube after the intermediate heating step to sizing hot rolling at a temperature equal to or greater than (Ar3+100°C), where Ar3 is the ferrite transformation start temperature calculated from formula (B) lower, and ending the hot rolling at a temperature of (Ar3+50°C) or higher; a direct tempering step of directly tempering the raw steel tube continuously from the second hot rolling step, the tempered raw steel tube being at a temperature equal to or greater than (Ar3+10°C) under conditions where the average cooling rate is 40°C / s more, and the cooling interruption temperature is 200°C or less; and a heat treatment step of subjecting the raw steel tube after the direct tempering step to at least one run of a heat treatment that tempers the raw steel tube after reheating to a temperature of 850 to 930°C, and subsequently tempers the raw steel tube by heating from 650 to 730°C, satisfying the recovery temperature Tr and the waiting time tW in the intermediate heating step according to the following formula (1). Ms = 545 - 330 χ (%C) - 7 χ (%S¡) - 23 χ (%Mn) - 14 χ (%Cr) - 5 χ (%Mo) + 2 χ (%AI) - 13 χ (%Cu) - 4 χ (%Nb) + 4 χ (%V) + 3 χ (%T¡) ... (A) Ar3 = 910 - 273 χ (%C) - 74 χ (%Mn)- 56 χ (%N¡) - 16 χ (%Cr) - 9 χ (%Mo) - 5 χ (%Cu)... (B) (Tr-Ms) < 10 + 0.0024 χ (tW)2... (1) In formulas (A) and (B), the atomic symbol represents the content of the element in % mass, and the content is (0) for the elements that are not contained. nrnonn / zznz / B / YiAi In the present invention, the steelmaking process is not particularly limited. For example, molten steel of the above composition can be produced using a known steelmaking process, such as using a converter, an electric furnace, or a vacuum melting furnace. For cost considerations, the molten steel is preferably melted by continuous casting. In continuous casting, the molten steel can be continuously melted into a common casting having a rectangular cross-section, such as a plate or a semi-finished metal casting, or it can be continuously melted directly into a casting having a circular cross-section, which is more suitable for hot rolling into a seamless steel tube.In the case of continuous casting into a casting that has a rectangular cross-section, the casting having a rectangular cross-section is heated to a predetermined heating temperature, and is not hot-rolled into a steel tube material that has a circular cross-section. The following describes a hot forming process for a seamless steel tube of a predetermined shape using steel tube material obtained after rolling a billet or heat treating a casting. In the present invention, the temperatures, including the steel tube material and raw steel tube temperatures, hot rolling temperature, cooling start temperature, cooling stop temperature, and heat treatment temperature, are surface temperatures of materials such as steel tube material and raw steel tube (the external surface of a tube in the case of a raw steel tube). These temperatures can be measured using a radiation thermometer or similar device. Heating Step of Steel Pipe Material Heating temperature: 1,200 to 1,300°C To form a seamless steel tube of a predetermined shape by hot rolling, steel pipe material is heated to the austenitic phase region of the steel. When the heating temperature of the steel pipe material is below 1,200°C, severe internal defects occur during punching, and defects detected in non-destructive testing after the final heat treatment of the steel pipe may not be satisfactory even after repair. From a defect prevention standpoint, the heating temperature of the steel pipe material is 1,200°C or higher. When the heating temperature of the steel pipe material is above 1,300°C, severe thickening of the austenite grains occurs in the steel. The impact of this thickening remains even after subsequent hot rolling, quenching, and heat treatment processes, and causes deterioration of the Kilolimit value.The upper limit of the heating temperature of the steel pipe material is therefore 1,300°C. First Step of Hot Rolling of Steel Tube (Drilling and Elongation Rolling Step) Final Lamination Temperature: 800°C or more In the first hot rolling of a seamless steel tube, the process begins with the punching roll, followed by the lengthening roll. When the temperature of the raw steel tube at the end of the lengthening roll is below 800°C, the high-temperature ductility of the steel decreases, and defects occur on the outer surface during hot rolling. This has adverse effects on the transformation behavior of the steel during the intermediate cooling described below and causes a deterioration of the Kilolimit value. For this reason, the final rolling temperature of the first hot rolling is 800°C or higher. The starting temperature for the first hot roll is not particularly limited. However, to prevent austenite grain thickening, the starting temperature for the first hot roll is preferably 1,280°C or lower. To prevent surface defects during hot rolling, the starting temperature for the first hot roll is preferably 1,150°C or higher. Intermediate Cooling Step of the Raw Steel Tube Cooling Start Temperature: 700°C or higher Intermediate cooling, when properly performed after lengthening in the first hot rolling, allows the steel tube to undergo bainite transformation, and the reverse transformation occurs during the subsequent intercooling. This greatly improves the Kilolimit value. When intercooling begins at a temperature below 700°C, the steel undergoes ferrite transformation before intercooling, and the reverse transformation behavior of the steel during the subsequent intercooling is adversely affected. This leads to a deterioration of the Kilolimit value. The starting temperature for intercooling is therefore 700°C or higher. Average Cooling Rate: 40°C / s more To allow bainite transformation in the raw steel tube, the average cooling rate of the intermediate quenching is 40°C / s or higher. As used herein, “average cooling rate” means the average cooling rate at the outer surface of the raw steel tube over a temperature range of 700°C to (Ms + 150°C), where Ms is the martensitic transformation onset temperature calculated using the following formula (A). With an average cooling rate of less than 40°C / s, it is not possible to initiate bainite transformation through the wall thickness of the raw steel tube. In this case, a region without bainite transformation exhibits the same transformation behavior as in the ordinary DQ-QT process, and the Kilolimit value cannot be improved.For this reason, the average cooling rate of the intermediate cooling is 40°C / s more, preferably 50°C / s more. The upper limit of the average cooling rate is not particularly restricted. However, the average cooling rate is preferably 100°C / s or less because it is extremely difficult at excessively high cooling rates to control the recovery temperature of the cooled raw steel tube (described later) within the predetermined temperature region. The method of cooling the raw steel tube is not particularly limited. However, it is preferable to cool the raw steel tube by water bath or mist application to the external surface of the tube so that intermediate cooling can be carried out after the raw steel tube is discharged from the hot rolling mill and enters the intermediate heating furnace, and thus the recovery temperature of the cooled raw steel tube can be more easily controlled within the predetermined temperature range. Recovery Temperature Tr: (Ms+150°C) or less ncnonn / zznz / B / YiAi For the bainite transformation of the raw steel tube, the recovery temperature Tr of the raw steel tube immediately after intermediate cooling needs to be (Ms+150°C) or less (Ms is the martensitic transformation temperature of the steel) so that the bainite transformation starts at least through the wall thickness of the raw steel tube. Figure 6 is a diagram representing time-dependent temperature changes at the outer surface, center of the wall thickness, and inner surface of a raw steel tube, as measured by heat transfer calculations. The raw steel tube is 28 mm thick and was cooled from 800°C. For the calculations, the raw steel tube was cooled by immersing the outer surface in water. The outer surface of the raw steel tube recovers after a transient temperature drop. The recovery temperature then converges to approximately the same temperatures measured at the center of the wall thickness and on the inner surface.This means that the temperature at the center of the wall thickness and the temperature on the inner surface of the steel pipe material have decreased to the same temperature range as the outer surface temperature when the recovery temperature on the outer surface of the steel pipe material has decreased to the predetermined temperature range. The Kilolimit value cannot reach its target value of 23.0 MPaT / m (Figure 7) when the recovery temperature Tr is above (Ms+150°C). The recovery temperature Tr is therefore (Ms+150°C) or lower, preferably (Ms+100°C) or lower, and most preferably (Ms+80°C) or lower. The onset temperature of the martensitic transformation Ms can be calculated from the following formula (A). Ms = 545 - 330 x (%C) - 7 x (%Si) - 23 x (%Mn) - 14 x (%Cr) - 5 x (%Mo) + 2 x (%AI) - 13 x (%Cu) - 4 x (%Nb) + 4 x (%V) + 3 x (%T¡) ... (A) In formula (A), the atomic symbol represents the element content in % by mass, and the content is zero (0) for elements that are not contained. The lower limit of the recovery temperature Tr is not particularly restricted. However, from an economic standpoint, the recovery temperature Tr is preferably equal to or greater than the martensitic transformation start temperature (Ms) because the fuel consumption rate in the subsequent intermediate heating step increases as the recovery temperature Tr decreases. The recovery temperature Tr is more preferably equal to or greater than (Ms + 30°C). It should be noted here that the Kilolimit value can still reach the target value of 23.0 MPaVm even if the recovery temperature Tr actually becomes equal to or less than the martensitic transformation start temperature (Ms). Intermediate Heating Step of the Raw Steel Tube Waiting Time tW before the Start of Intermediate Heating As discussed previously, the cooling interruption temperature of the intermediate cooling step (specifically, the recovery temperature after intermediate cooling) and the time before the start of the subsequent intermediate heating step are important. The inventors of the present device found that the recovery temperature Tr immediately after intermediate cooling and the waiting time tW before the start of intermediate heating have combinations with which the Kilolimit value can reach the target value of 23.0 MPa x 10⁻¹. Specifically, the waiting time tW before the start of intermediate heating needs to be longer for higher recovery temperatures Tr. Conversely, shorter waiting times tW are sufficient for lower recovery temperatures Tr.Referring to Figure 7, the inventors of the present obtained formula (1) by approximating a quadratic curve for the limit of the target Kilílimit value, using the recovery temperatures Tr and waiting times tW obtained in a simulation experiment. (Tr-Ms) < 10 + 0.0024 χ (tW)2... (1) When the value of (Tr-Ms) is smaller than the value on the right-hand side of formula (1), the bainite transformation can proceed almost completely until the intermediate heating begins, and the reverse transformation can take place during the subsequent intermediate heating, allowing the Kilolimit value to reach the target value of 23.0 MPaum through grain refining. From a production efficiency standpoint, the waiting time tW before the start of intermediate heating is 300 seconds or less, preferably 250 seconds or less, and most preferably 200 seconds or less. The lower limit of the waiting time tW before the start of intermediate heating is not particularly restricted.However, considering the restrictions of the equipment used for intermediate cooling to intermediate heating processes, the waiting time tW is preferably 30 seconds or more, more preferably 100 seconds or more, provided that formula (1) is satisfied. Intermediate Heating Temperature: 800 to 1,000°C Intermediate heating is performed to promote grain refinement through the reverse transformation of the intermediate-cooled steel tube and to apply supplementary heat to the steel tube during the dimensioning of a seamless steel tube. When the intermediate heating temperature is below 800°C, the steel tube continues to undergo reverse transformation, and the grains are not refined as intended. Because this leads to a decrease in the Kilolimit value, the intermediate heating temperature is 800°C or higher. The intermediate heating temperature is 1,000°C or lower because severe oiling, rather than refinement, of the grains occurs as a result of grain growth when the intermediate heating temperature is above 1,000°C. Second Hot Rolling Step of Steel Tube (Sizing Rolling Step) Lamination Start Temperature: (Ar3+100°C) or higher The intermediate heating is followed by sizing rolling (second hot rolling), a final hot rolling step. Rolling causes grain mixing in the microstructure and decreases the Kilolimit value when the sizing rolling start temperature is lower than (Ar3 + 100°C), where Ar3 is the ferrite transformation start temperature. For this reason, the second hot rolling start temperature is (Ar3 + 100°C) or higher. The ferrite transformation start temperature (Ar3) can be calculated from the following formula (B). Ar3 = 910 - 273 χ (%C) - 74 χ (%Mn) - 56 χ (%N¡) - 16 χ (%Cr) - 9 χ (%Mo) - 5 χ (%Cu)... (B) In formula (B), the atomic symbol represents the element content in % by mass, and the content is zero (0) for elements that are not contained. The upper limit of the starting temperature for the second hot rolling stage is not particularly restricted. However, from the standpoint of production efficiency, the starting temperature for the second hot rolling stage is preferably 1,000°C or lower. Final Rolling Temperature: (Ar3+50°C) or more The final rolling temperature of the second hot rolling is (Ar3+50°C) or higher because rolling causes grain mixing in the microstructure, and decreases the Kilolimit value when the final sizing rolling temperature is less than (Ar3+50°C). The upper limit of the final rolling temperature of the second hot rolling is not particularly limited, and is preferably 950°C or less. Direct Tempering Step Direct Tempering Starting Temperature: (Ar3+10°C) or more The sizing roll (second hot roll) is followed by direct quenching (DQ) of the raw steel tube. When the start temperature of the direct quenching is lower than (Ar3+10°C), ferrite transformation occurs during the quenching process, and the effect of the direct quenching becomes insufficient. For this reason, the start temperature of the direct quenching is (Ar3+10°C) or higher. The upper limit of the starting temperature for the direct tempering step is not particularly restricted, and is preferably 900°C or less. Average Cooling Rate: 40°C / s more The average cooling rate for direct quenching is 40°C / s or more because the effect of direct quenching becomes insufficient when the average cooling rate is less than 40°C / s. The average cooling rate for direct quenching is preferably 50°C / s or more. As used herein, “average cooling rate” means the average cooling rate at the outer surface of the as-finished steel tube over a temperature range of (Ar3+10°C) to 200°C. The upper limit of the average cooling rate is not particularly restricted. However, from the standpoint of preventing work-hardening cracking during cooling, the average cooling rate is 100°C / s or less. Cooling Interruption Temperature: 200°C or less The cooling cutoff temperature for direct tempering is 200°C or less because the cooling effect of direct tempering becomes insufficient when the cooling cutoff temperature is above 200°C. The cooling cutoff temperature for direct tempering is preferably 150°C or less, and more preferably 100°C or less. The lower limit of the cooling interruption temperature is not particularly restricted. However, from the standpoint of cooling efficiency, the cooling interruption temperature is preferably at least room temperature, and more preferably 50°C or higher. The cooling method in direct quenching is not particularly restricted, and cooling can be achieved, for example, by immersing the raw steel tube in a water tank, by spraying water from the inside and outside of the raw steel tube, or by applying mist. Any of these methods may be used, provided the specified average cooling rate can be achieved. Heat Treatment Step Tempering Reheating Temperature: 850 to 930°C ncnonn / zznz / B / YiAi The direct tempering step is followed by quenching, which involves reheating the raw steel tube to adjust its strength to API grade C110. When the quenching reheat temperature is below 850°C, the austenite transformation of the raw steel tube does not proceed completely, and the untransformed region causes a decrease in strength. For this reason, the quenching reheat temperature is 850°C or higher, preferably 870°C or higher. When the quenching reheat temperature is above 930°C, grain coarsening occurs, and the Kilolimit value decreases. For this reason, the quenching reheat temperature is 930°C or lower, preferably 900°C or lower. The cooling method in reheat tempering is not particularly limited, as is the case with direct tempering. For example, cooling can be achieved using any method, including immersing the raw steel tube in a water tank, spraying the inside and outside of the raw steel tube with water, and applying mist. Tempering temperature: 650 to 730°C Hardening by reheating is followed by tempering to adjust the raw steel tube to a strength equivalent to API C110 standards. When the tempering temperature is below 650°C, the strength of the steel tube increases excessively, and the Kilolimit value decreases. For this reason, the tempering temperature is 650°C or higher, preferably 680°C or higher. When the tempering temperature is above 730°C, the reverse transformation occurs in parts of the steel, and the strength decreases significantly. For this reason, the tempering temperature is 730°C or lower, preferably 710°C or lower. The quenching and reheating (QT) process is performed at least once. The quenching and reheating process can be performed two or more times to obtain even higher Kilolimit values. Examples The present invention is described in greater detail below by way of Examples. It should be noted that the present invention is not limited by the following Examples. In the steels of the compositions shown in Table 2, steels A, B, and C were produced using a converter steelmaking process and cast into semi-finished castings by continuous casting. In Table 2, the symbol ≡ indicates that the element was not intentionally added, meaning that the element may be absent (0%) or may be accidentally present. The semi-finished casting was hot-rolled into steel tube stock having a circular cross-section, and the steel tube stock was machined to produce a block for the hot-rolling experiment. For the other steels, blocks for the hot-rolling experiment were produced using a vacuum melting furnace.These were subjected to hot plate rolling conducted as a simulation of the hot rolling, intercooling, interheating, hot rolling, and direct tempering of a seamless steel tube, using a small-scale rolling mill, a cooling device, and a heating furnace. The plate thicknesses of the rolled materials, and the heating, rolling, and cooling conditions are as shown in Table 3-1 and Table 3-2. The plate temperature of the rolled material was measured with a thermocouple embedded in the surface on one side of the rolled material. The hot-rolled steel plates were then subjected to a heat treatment of quenching and tempering using the reheating conditions shown in Table 3-1 and Table 3-2. From the heat-treated material, a JIS 14A round bar tensile test specimen was taken in accordance with JIS Z2241 (2011). The test specimen was used for a common temperature tensile test performed according to JIS Z2241, and the creep (YS) of the heat-treated material was measured. To confirm grain refinement, a microscopic sample was taken from the same heat-treated material. The sample was polished to a mirror finish and etched with a picral solution (a picric acid-ethanol mixture). After revealing the pre-austenitic grain boundary, micrographs of four randomly selected fields were taken using a light microscope at 1000x magnification. The pre-austenitic grain size number (PEN) of the photographed PEN was then measured using the y-intercept method in accordance with JIS G0551 (2013). The pre-austenitic grain size (PEN) was measured as a grain size number in accordance with ASTM E112. For the Kilolimit evaluation, a DCB test specimen measuring 9.5 mm thick, 25.4 mm wide, and 101.6 mm long was taken according to NACE TM0177 Method D. A total of nine DCB test specimens were taken from each sample and subjected to a DCB test. The DCB test was conducted in a test bath containing an aqueous solution at 24°C of 5 wt% NaCl and 0.5 wt% CH3COOH saturated with 1 atm (0.1 MPa) of hydrogen sulfide gas. After inserting a wedge, the DCB test specimen was immersed in the test bath for 336 hours under predetermined conditions, and the length of a crack generated in the DCB test specimen while immersed in the solution was measured. The open wedge stress P was also measured on the specimen. Then the Kissc (MPa / m) was calculated using the following formula (0). nrnonn / zznz / e / YiAi In formula (0), h is the arm height (height of each arm) of the DCB test specimen, B is the thickness of the DCB test specimen, and Bn is the net thickness of the DCB test specimen. These values ​​are specified in NACE TM0177 Method D. Based on the expected maximum notch defect and load application conditions for oilfield tubular products, the target Kilolimit value was established as 23.0 MPaam or higher. For the Kilolimit value calculation, the wedge was used in three different thicknesses: 2.76 mm, 2.89 mm, and 3.02 mm, and each was used for at least three test specimens. A Kilolimit value was calculated following the procedures described with reference to Figure 1, using the calculated Kissc values. The yield strengths and ultimate tensile strengths (Kilometer) of the heat-treated materials are presented in Table 4-1 and Table 4-2. The yield strength falls within the range of the present invention when it is 758 MPa or greater. The ultimate tensile strength (Kilometer) falls within the range of the present invention when it is 23.0 MPa or greater. The ultimate tensile strength (Kilometer) is preferably 24.0 MPa or greater, and more preferably 25.0 MPa or greater. Table 2 No. of Maple Composition (% in mass) C Si Mn PS Cr Mo Al Cu Nb VB Ti ON Ca A 0.24 0.03 0.59 0.006 0.0005 0.96 0.66 0.068 0.06 0.033 0.038 0.0019 0.001 0.0009 0.0025 - B 0.26 0.04 0.56 0.005 0.0006 1.04 0.73 0.066 0.05 0.031 0.032 0.0022 0.002 0.0010 0.0028 0.0011 C 0.25 0.19 0.51 0.008 0.0007 0.92 0.80 0.067 0.07 0.026 0.027 0.0027 0.003 0.0012 0.0037 - D 0.26 0.11 0.65 0.007 0.0008 1.09 0.62 0.052 0.03 0.044 0.025 0.0016 0.003 0.0014 0.0035 0.0014 E 0.23 0.34 0.69 0.009 0.0009 1.18 0.52 0.078 0.02 0.049 0.021 0.0015 0.004 0.0018 0.0044 - F 0.27 0.01 0.46 0.010 0.0007 0.81 0.89 0.041 0.08 0.022 0.048 0.0029 0.005 0.0009 0.0027 0.0017 G 0.21 0.30 0.68 0.010 0.0009 1.19 0.88 0.077 0.02 0.048 0.024 0.0022 0.002 0.0008 0.0026 - H 0.24 0.33 0.43 0.009 0.0010 1.15 0.89 0.079 0.03 0.047 0.022 0.0019 0.003 0.0009 0.0029 - I 0.23 0.34 0.68 0.010 0.0010 0.74 0.88 0.078 0.07 0.047 0.023 0.0017 0.002 0.0011 0.0031 - J 0.24 0.35 0.67 0.010 0.0009 1.17 0.39 0.077 0.05 0.046 0.024 0.0018 0.002 0.0010 0.0028 - K 0.23 0.35 0.66 0.009 0.0008 1.18 0.87 0.078 0.02 0.048 0.022 0.0011 0.003 0.0009 0.0033 - L 0.27 0.02 0.47 0.009 0.0008 0.82 0.88 0.043 0.09 0.020 0.050 0.0028 0.005 0.0016 0.0047 0.0002. ncnQnn / zznz / e / Ywi Tabla 3-1 seuoiOB / uesqo UJ LU UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ UJ LU UJ O LU OO Tratamiento térmico de recalentamiento CXJ O • 8 α • • • • (Co) SO CO CO co • • • I--- o__ 8 8 co 8 g LO 8 sg 05 8 co 05 co LO g O 05 LO CXJ § 8 8 LO 8 co 8 σ 8 LO co co 8 co co 8 co 8 8 § 8 8 8 8 8 8 8 05 8 8 LO co co 8 8 8 8 8 co 8 OQ Tem Pfinal (°C) co co §5 s co LO co 8 05 05 5 co co LO co co co S io LO 55 LO LO 8 ¡o Velocidad de enfriamient o promedio (°C / s) s LO LO LO s co co LO lo 8 LO LO co LO 8 LO LO LO 5 LO LO co LO LO LO S 8 LO LO 8 Temp. initial (°C) 8 g LO co co LO co ss co co CO co 05 co co LO co 8 8 LO co 8 CXJ co 8 co co 8 8 g Segunda laminación en caliente Temp, final (“O CXJ co CE» co LO LO co co co co co Cn CXJ co co co co co co co 8 co LO co co LO CXJ co 8 co CXJ co 8 05 05 03 CO 8 05 CO co CXJ co co OT co Temp.initial (°C) co co 8 co 8 LO co co 8 co cn 03 LO co co co LO co 8 co § co LO co co 8 LO LO 03 LO co 8 8 co LO co co LO co co 8 Value on the right side of formula (1) co LO Q co Q oo> co co LO co LO 8 05 no 8 8 co rS CO CXJ LO 8 s Tr-Ms cxj LO co LO lo r- LO 8 CXI co ss ES S co CXJ ES 03 8 ES δ CXJ 8 8 Intermediate heating Temp. surface (“O 8 8 8 8 8 Sf S 05 LO 05 8 8 co 8 s 05 8 co 8 8 s 8 s 8 8 s tw (sec) CXJ co en co §5 05 o co s LO LO CXJ 3 CXJ LO CXJ LO LO CXJ 8 co CXJ CXJ ¡o CXJ CXJ co CXJ 8 LO CXJ 8 co ES s Intermediate cooling h= υ o Es co co gg en 8 cn 05 co CXJ LO LO LO 8 LO 8 05 05 s co ES O 8 CXJ co CXJ Average cooling rate (°C / s) CXJ S 8 s LO co s LO LO LO LO 8 LO LO 05 CXJ 05 LO 5 O) CO LO LO LO LO 8 LO 8 Initial temp. (°C) oon 1010 1015 1041 1051 1047 &> co 05 LO ORL 1035 1025 1040 8 8 1040 1095 1020 1020 1025 1021 1033 1041 1022 First hot rolling Temp.final (°C) 1129 1043 1050 1100 1104 1103 8 cn CO 05 1140 1068 1050 1085 8 CXJ co 1067 1120 1040 1044 1060 1060 1050 Temp initial <°C) 1200 1155 1160 1180 1180 1180 OSSf 1220 1240 1210 1160 1160 1150 1000 1000 1155 1200.61 o 1161 1 1160 Temp. of Heats ment. (°C) 1250 1210 1210 1230 1230 1290 1250 1210 0661 1225 0661 (Do) SJV 05 σ> 05 co CXJ co CXJ co co co CO CXJ 8 CXJ CXJ CXJ co co LO 8 CXJ CXJ 8 o o__ co LO co LO CXJ Es ES co CXJ co LO XJ 8 CXI 8 LO C 8 8 8 8 co CXI Thickness (mm) CM co CXJ co CXJ co LO CXJ co CXJ co CXJ co CXJ LO CXJ LO CXJ LO CXJ LO CXJ 8 CXJ LO CXJ LO CXI LO CXJ No. LO CXJ LO CXJ 8 of sample with mmm OO with δ CXJ or lZj UJ THIS LO UJ CXI <5 í 2= —5 Zj No. of Steel < with co OOO d UJ UJ UJ UJ UJ U_ O ΖΓ — —5 ________1. *3 Ar3 = 910 - 273 x (%C) - 74 x (%Mn) - 56 x (%N¡) -16 x (%Cr) - 9 x (%Mo) - 5 x (%Cu) '4 (Tr-Ms) < 10 + 0.0024 Example... x (tW) Comparative Example S > XC IX IX CC α c Ca c Table 3-2 seuoiOE / ussqo LJJ EYE EYE Ljj EYE o EYE EYE o UJ ion x (%Mo) + 2 x (% Al) -13 x (%Cu) - 4 x (%N b) + 4 x (%-9) + (%Ti) + 3% x (% x (%Cu) Reheat heat treatment (Oo) 21 (Oo) 20 (Oo) ti S g LO fe δ o3 fe fe oo 05 fe fe co 05 fe fe fe fe fe fe fe fe fe fe fe fe co si O5| with faith faith with OQ Temp. final (°C) s LO 5 LO <s g lo cm 05 oo o co fe δ velocidad de enfriamiento promedio (°c s) o>δΐ LO LO LO δ LO faith Temp. initial (°C) OO So LO CO co co g δ fe LO OO OO s fe fe ss LO OO OO co co co co co co co Second hot rolling Temp. final (°C) GO CM co co oo δ co co co fe ση δ δ co co 05 05 col col CO δ δ oo co oo s δ co Temp. initial (°Q co fe oo CM fe co fe oí co co fe co fe co fe co 1030 co co co oo fe oo co co S co s co fe co fe co Value on the right side of the formula (1) CM f2 LO g| coi S 05 co fe δ <5 co 05 LO CM oo LO co 05 co co 05 co Tr-Ms LO fe LO δΐ co SI OJ LO CO δ δ fe CO Yes fe CO CO fe s Intermediate heating Surface temp. (°C) S CM σ> fe § CM 05 fe SSS fe 1056 | SI fe ss SS 05 s fe tw (sec g CM δ δ 05 LO r-· co 05 fe δ δ s co r^· δ co fe 05 Intermediate cooling O co COI Lol oo loI g fe LO δ OO 05 05 s CO 05 cm g LO CM LO o Average cooling rate (°C / s) LO co LO LO LO LO fe LO LO oo co LO LO col LO LO LO δ LO LO LO LO LO LO fe fe co LO Temp.initial (°C) 1098 ¿col CO 05 1093 1027 CM CM 05 1144 δ δ col 1100 1100 1099 0011 1091 1100 1094 1100 1094 1100 1094 1100 1094 1100 001 Prime lamination temp. final (°C) t£U 1084 en 1124 1075 CT> CM 05 co 1133 ej 1132 1130 1124 ten 1128 1134 1129 1132 1130 interval of the present invention >i) -23 x (4%) xr -51 -C ÓMn) - 56 x (%Ni) -16 x (%Cr) (i) Tem P· inicia I (°C) 0021 1160 gg LO LO 1225 1265 § 1000 1200 gg 1200 0021 fe gg 1200 g 1200 g Temp. of Heating ment. (°C) 1250 | 1230 | 1265 | 1250 | 1230 | 1269 | I 1310 I 1200 | 1200 | 1250 | 1250 | 1250 | 1250 | 1250 | 1250 | I 1250 I 1250 | 1250 | 1250 | 1250 | (Oo) w CM OO 05 CM CO 05 05 05 05 05 05 05 05 05 CD 05 05 05 05 Ms (CC) lo CM CM lo fe CO s lo S lo lo LO LO s lo S ss LO LO Ί The underline means outside the M ss (35% - 35 - 35 x -C) 7 x (%S '3 Ar3 = 910 - 273 x (%C) - 74 x (°z *4 (Tr-Ms)< 10 +0.0024 xlJW)2 CE: Ejemplo Comparativo Espesor (mm) co CO OO CM co CM C© co co co CM co CM co CM co CM co s co CM co CM oo CM co CM co co CM No.de Muestra 3 S δ mo co < ​​< LO < co 05 No. de Acero C CO O < CÜ O . <c <c <c < <cnmonn / z / nzE / YiAi Tabla 4-1 Steel No. Sample No. ASTM Number Previous Austenitic Grain Size YS (MPa) Kilolimit (MPa / m) Remarks A A1 11.0 808 24.3 Example of Present A A2 10.5 774 24.6 Example of Present A A3 12.5 824 25.2 Example of Present B B1 11.0 819 24.1 Example of Present B B2 11.0 803 24.2 Example of Present B B3 13.0 808 25.5 Example of Present C 01 11.0 782 23.8 Example of Present C 02 10.5 771 23.3 Example of Present C 03 11.0 778 23.6 Example of Present D D1 10.5 833 23.5 Example of the Present Tense D D2 10.5 766 23.9 Example of the Present Tense E E1 10.5 859 23.0 Example of the Present Tense E E2 10.5 759 23.3 Example of the Present Tense E E3 10.5 762 23.4 Example of the Present Tense E E4 10.5 764 23.3 Example of the Present Tense E E5 10.5 772 23.2 Example of the Present Tense F F1 10.5 839 23.4 Example of the Present Tense F F2 10.5 763 23.8 Example of the Present Tense G G1 10.5 669 25.8 Comparative Example H H1 10.5 744 24.2 Comparative Example I 11 10.5 751 24.3 Comparative Example J J1 10.5 725 24.9 Comparative Example K K1 10.5 696 25.2 Comparative Example L L1 11.0 822 22.1 Comparative Example. The underline means outside the range of the present invention ncnonn / zznz / e / YiAi Table 4-2 Steel No. ASTM Sample No. Previous Austenitic Grain Size Number YS (MPa) Kilolimit (MPa / m) Remarks A A4 9.5 799 21.5 Comparative Example B B4 10.0 812 22.2 Comparative Example C C4 10.0 766 22.0 Comparative Example A A5 10.0 803 22.8 Comparative Example B B5 10.0 809 21.8 Comparative Example C C5 9.5 762 21.4 Comparative Example A A6 10.0 803 22.7 Comparative Example A A7 10.0 804 22.6 Comparative Example A A8 10.0 805 22.5 Comparative Example A A9 9.5 791 22.3 Comparative Example A A10 9.5 794 22.4 Comparative Example A A11 9.0 760 21.9 Comparative Example A A12 10.0 804 22.8 Comparative Example A A13 10.0 807 22.5 Comparative Example A A14 9.5 795 22.2 Comparative Example A A15 10.0 806 22.5 Comparative Example A A16 9.0 759 21.4 Comparative Example A A17 11.0 741 24.9 Comparative Example A A18 11.0 728 25.3 Comparative Example A A19 11.0 914 22.1 Comparative Example The underline means outside the range of the present invention nrnonn / zznz / e / YiAi As shown in Tables 3-1 and 3-2 and in Tables 4-1 and 4-2, the creep was satisfactory, and the Kilolimit value was excellent in all the examples hereof (Samples Nos. A1 to A3, B1 to B3, C1 to C3, D1 to D2, E1 to E5, and F1 to F2) in which the steel compositions and manufacturing conditions satisfied the ranges of the present invention, and the value of (Tr-Ms) was equal to or less than the value on the right-hand side of the above formula (1), where Tr is the recovery temperature, and Ms is the temperature at which the martensitic transformation of the steel begins. In contrast, none of the samples Nos. G1, H1, 11, J1, and K1 from the Comparative Examples met the target creep. In Sample No. L1 of the Comparative Example, large numbers of coarse non-metallic oxide inclusions were observed, and the Kilolimit value did not meet the target value. In the Comparative Examples (samples Nos. A4, B4, and C4) where the steel compositions met the preferred ranges but the recovery temperature Tr after intermediate cooling exceeded (Ms+150°C), no bainite transformation occurred after intermediate cooling and before the start of intermediate heating. As a result, grain refinement was insufficient, and the Kilolimit value did not meet the target value. In the Comparative Examples (samples Nos. A5, B5, C4, and C5) where the (Tr-Ms) value exceeded the value on the right-hand side of formula (1), the bainite transformation began but did not end before the start of reheating. As a result, grain refining was insufficient, and the Kilolimit value did not meet the target value. Aggregation of the austenite grains occurred, and the Kilolimit value did not satisfy the target value in the Comparative Example (sample No. A6) in which the heating temperature of the steel tube material was above the upper limit of the present invention, and in the Comparative Example (sample No. A10) in which the intermediate heating temperature was above the lower limit of the present invention. In the Comparative Examples (Samples Nos. A7, A12) in which the final temperature of the first and second hot rolling was below the lower limit of the present invention, the low rolling temperatures had adverse effects on the transformation in the subsequent cooling process, and the Kilolimit value did not satisfy the target value. In the Comparative Example (sample No. A8) in which the starting temperature of the intermediate cooling after the first hot rolling was below the lower limit of the present invention, and in the Comparative Example (sample No. A13) in which the starting temperature of the direct quenching cooling was below the lower limit of the present invention, ferrite transformation occurred before the intermediate cooling (sample No. A8) and before the direct quenching (sample No. A13), and the transformed microstructure had grain mixing, with the result that the Kilolimit value did not satisfy the target value. In the Comparative Example (sample No. A9), where the average cooling rate of the intermediate cooling was below the lower limit of the present invention, no bainite transformation occurred after the intermediate cooling and subsequent recovery and before the start of reheating. As a result, no refining of the grains took place, and the Kilolimit value did not meet the target value. In the Comparative Example (sample No. A11) in which the intermediate heating temperature was below the lower limit of the present invention, the reverse transformation did not end at the time of reheating, and the refining of the grains did not take place, with the result that the Kilolimit value did not satisfy the target value. The effect of direct tempering was insufficient in the Comparative Example (Sample No. A14), in which the average cooling rate of direct tempering was below the lower limit of the present invention, and in the Comparative Example (Sample No. A15), in which the cooling interruption temperature of direct tempering was above the upper limit of the present invention. As a result, grain refining did not take place, and the Kilolimit value did not meet the target value. In the Comparative Example (sample No. A16) in which the heating temperature of the reheat quenching in the reheat heat treatment was above the upper limit of the present invention, thickening of the austenite grains occurred, and the Kilolimit value did not satisfy the target value. Conversely, in the Comparative Example (sample No. A17) in which the reheating temperature was below the lower limit of the present invention, some regions of the steel were left untransformed after quenching, and the creep did not meet the target value. In the Comparative Example (sample No. A18) in which the tempering temperature after reheating tempering was above the lower limit of the present invention, reverse transformation of the steel occurred during tempering, and the creep did not satisfy the target value. nrnonn / zznz / e / YiAi Conversely, in the Comparative Example (sample No. A19) in which the tempering temperature was below the lower limit of the present invention, the strength increased excessively, and the Kilolimit value did not satisfy the target value.< / c> < / s>

Claims

1. A high-strength seamless steel tube characterized in that it has a yield strength of 758 MPa or more, and a Kilolimit value of 23.0 MPa or more according to a sulfide stress corrosion cracking resistance evaluation index, wherein Kilolimit is a value determined from the intersection between (i) a linear regression line created by a stress intensity factor Kissc obtained from a DCB (Double Cantilever Beam) test performed multiple times under different test conditions, and a stress intensity factor Kiapiked at the tip of a notch on a test specimen prior to the start of the DCB test, and (ii) a straight line on which Kissc and Kiapiked are one to one.

2. The high-strength seamless steel tube according to claim 1, further characterized in that the microstructure of the steel with a previous austenitic grain size of 10.5 or more in terms of a grain size number complies with ASTM E112.

3. The high-strength seamless steel tube according to claim 1 or 2, further characterized in that it has a composition comprising, in % by mass, C: 0.23 to 0.27%, Si: 0.35% or less, Mn: 0.45 to 0.70%, P: 0.010% or less, S: 0.0010% or less, Cr: 0.80 to 1.20%, Mo: 0.50 to 0.90%, Al: 0.080% or less, Cu: 0.09% or less, Nb: 0.050% or less, V: 0.050% or less, B: 0.0015 to 0.0030%, Ti: 0.005% or less, O: 0.0020% or less, and N: 0.0050% or less, in the specification in which the remainder is Fe and accidental impurities.

4. The high-strength seamless steel tube according to claim 3, further characterized in that the composition additionally comprises, in % by mass, Ca: 0.0020% or less.

5. A method for manufacturing high-strength seamless steel tubing according to any one of claims 1 to 4, the method being characterized in that it comprises: a heating step of a steel tubing material to a heating temperature in a temperature region of 1,200 to 1,300°C; a first hot rolling step of hot rolling the steel tubing material by piercing and elongating the steel tubing material with a final rolling temperature of 800°C or more;an intermediate cooling step of cooling a raw steel tube after the first hot rolling step, the raw steel tube being cooled from a cooling start temperature of 700°C or more under conditions where the average cooling rate is 40°C / s or more, and the recovery temperature Tr of the raw steel tube at the tube surface is (Ms+150°C) or less, where Ms is a martensitic transformation start temperature; an intermediate heating step of heating the raw steel tube after the intermediate cooling step, the raw steel tube being heated to a surface temperature of 800 to 1,000°C after a holding time tW of 300 seconds or less being charged into a reheating furnace;a second hot rolling step of subjecting the raw steel tube after the intermediate heating step to sizing hot rolling at a temperature equal to or greater than (Ar3+100°C), where Ar3 is a ferrite transformation start temperature, and hot rolling is completed at a temperature of (Ar3+50°C) or higher; a direct quenching step of directly quenching the raw steel tube continuously from the second hot rolling step, the raw steel tube being quenched from a temperature equal to or greater than (Ar3+10°C) under conditions where the average cooling rate is 40°C / s or higher, and the cooling interruption temperature is 200°C or lower;and a heat treatment step of subjecting the raw steel tube after the direct tempering step to at least one run of a heat treatment that tempers the raw steel tube after reheating to a temperature of 850 to 930°C, and continuously tempers the raw steel tube by heating from 650 to 730°C, 10 satisfying the recovery temperature Tr and the holding time tW in the intermediate heating step the following formula (1): (Tr-Ms) < 10 + 0.0024 x (tW)2 (1).;