Pipeline steel and manufacturing method therefor

Abstract

The present invention relates to a pipeline steel and a manufacturing method therefor. In addition to Fe and inevitable impurities, the pipeline steel further contains the following chemical elements in percentage by mass: 0.015-0.040% of C, 0.10-0.30% of Si, 1.18-1.45% of Mn, 0<P≤0.008%, 0<S≤0.0010%, 0.05-0.35% of Cu, 0.05-0.30% of Ni, 0.15-0.30% of Cr, 0.02-0.09% of Mo, 0.025-0.055% of Nb, 0.005-0.020% of Ti, 0.0010-0.0040% of Ca, 0.010-0.040% of Alt, 0<B≤0.0004%, 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, 0<Sn≤0.0050%, 0<Sb≤0.0025%, 0<Bi≤0.0030%, 0<Pb≤0.0050%, and 0<As≤0.0030%, and the chemical elements satisfy: Nb+Ti≤0.07%; Ca / S≥2; Cu+Ni+Cr+Mo≤0.80%; 0<N+O+H≤0.0070%; and 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0; a hydrogen induced cracking (HIC) sensitivity coefficient Ks≤0.23, and Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19. The obtained pipeline steel can reach an X60 / X65 strength level, and has good low-temperature toughness, and excellent HIC resistance and SSC resistance.

Description

Pipeline steel and manufacturing method thereof Technical Field

[0001] The present invention belongs to the field of metal materials and their manufacturing, and in particular relates to pipeline steel and a manufacturing method thereof. Background Art

[0002] With the continuous development of conventional oil and gas fields, reserves are constantly decreasing, and the proportion of H2S-containing sour oil and gas fields is further increasing. This requires that pipelines have acid resistance. Designers of different acid-resistant pipeline projects will develop personalized acid resistance indicators based on the pipeline's service conditions, including NACE standards for hydrogen-induced cracking (HIC) resistance and hydrogen sulfide stress corrosion cracking (SSC) resistance, as well as strict low-magnification rating and flaw detection requirements.

[0003] Chinese patent CN202010552764.7 discloses "A magnesium-containing X65 pipeline steel with excellent acid resistance and its production method." The pipeline steel's composition and weight percentages are as follows: C: 0.03-0.05%, Si: 0.17-0.20%, Mn: 0.40-0.70%, Ni: 0.12-0.15%, Cr: 0.45-0.48%, Mo: 0.10-0.15%, Mg: 0.002-0.004%, Nb: 0.06-0.09%, Ti: 0.01-0.02%, P ≤ 0.008%, S ≤ 0.010%, with the remainder being Fe and unavoidable impurities. This patent uses a Mg treatment method to remove inclusions, resulting in a low Mn content in the chemical composition. Therefore, by adding a large amount of Nb, Cr, and a certain amount of Mo to ensure strength, the acid-resistant pipeline steel with X65 strength is achieved. However, the use of Mg to treat inclusions makes it difficult to control the stable addition, and the control effect on inclusions such as MnS is poor, which has an adverse effect on acid resistance and toughness. The obtained acid-resistant pipeline steel has a Charpy impact energy of ≥200J at -20℃ and a DWTT SA% of ≥90% at -15℃.

[0004] Chinese patent CN201510870777.8 discloses a "thick gauge acid-resistant X60MS hot-rolled coil and its manufacturing method." The composition of the steel plate is as follows by weight: C: 0.050-0.070%, Si: 0.15-0.25%, Mn: 1.35-1.45%, Nb: 0.025-0.045%, Ti: 0.008-0.025%, Mo: 0.06-0.10%, Cr: 0.08-0.18%, Als: 0.015-0.045%, P≤0.015%, S≤0.0015%, N≤0.006%. Since an increase in C content significantly increases the degree of segregation in the steel, increases the HIC sensitivity of the steel, and reduces the low-temperature toughness of the steel to a certain extent, the resulting steel has a -40°C Charpy impact energy ≥300J and a -30°C DWTT. SA%≥85%.

[0005] Chinese patent CN202011197059.6 discloses "A X60 / X65 grade acid-resistant pipeline steel and its preparation method", whose chemical composition and mass percentage are as follows: C: 0.02~0.06%, Si: 0.10~0.20%, Mn≤0.90%, Alt: 0.020~0.050%, Nb: 0.060~0.080%, Ti≤0.020%, Cr≤0.30%, Ni≤0.20%, Cu≤0.40%, P≤0.012%, S≤0.002%, N≤0.0060%, B≤0.0005%, Mn+Cr≤1.2%, and the rest are iron and unavoidable impurity elements. The design of the patented alloy composition adds a large amount of Nb (0.060-0.080%). The high Nb content increases the tendency of Nb-containing particles in the steel to gather in the core, and the probability of HIC cracks appearing in the core of the steel plate increases accordingly.

[0006] It can be found from existing patents that the content of segregation-prone elements such as C and Nb added to existing acid-resistant pipeline steels is too high, which easily leads to segregation and inclusions in the core, increases the steel's sensitivity to HIC, increases the possibility of hydrogen-induced cracking, and makes it difficult to further improve the steel's low-temperature toughness.

[0007] Summary of the Invention

[0008] The object of the present invention is to provide a high-toughness X60 / X65 grade acid-resistant pipeline steel and a manufacturing method thereof, wherein the pipeline steel can reach the X60 / X65 strength level and has good low-temperature toughness and excellent acid resistance (such as HIC resistance and SSC resistance). t0.5The yield strength is 430~570MPa, the tensile strength is 520~690MPa, the yield strength ratio ≤0.92, the elongation A50 ≥30%, the hardness HV10 ≤220, the Charpy impact energy at -49℃ ≥400J, and the DWTT SA% at -35℃ ≥90%.

[0009] To achieve the above object, the present invention provides a pipeline steel, which, in addition to Fe and unavoidable impurities, further contains the following chemical elements by mass percentage: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0<P≤0.008%, 0<S≤0.0010%, Cu: 0.05-0.35%, Ni: 0.05-0.30%, Cr: 0.15-0.30%, Mo: 0.02-0.09%, N b: 0.025-0.055%, Ti: 0.005-0.020%, Ca: 0.0010-0.0040%, Alt: 0.010-0.040%, 0<B≤0.0004%, 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, 0<Sn≤0.0050%, 0<Sb≤0.0025%, 0<Bi≤0.0030%, 0<Pb≤0.0050%, 0<As≤0.0030%, and satisfy the following:

[0010] Nb+Ti≤0.07%;

[0011] Ca / S≥2;

[0012] Cu+Ni+Cr+Mo≤0.80%;

[0013] 0<N+O+H≤0.0070%;

[0014] 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0; and

[0015] The hydrogen-induced cracking sensitivity coefficient Ks≤0.23, Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19, where the element symbol is substituted into the value before the mass percentage sign of the corresponding element content.

[0016] In this article, “inevitable impurities” refer to other inevitable impurities except P and S.

[0017] Preferably, the pipeline steel contains the following chemical elements by mass percentage: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0<P≤0.008%, 0<S≤0.0010%, Cu: 0.05-0.35%, Ni: 0.05~0.30%, Cr: 0.15~0.30%, Mo: 0.02~0.09%, Nb: 0.025~0.055%, Ti: 0.005~0.020%, Ca: 0.0010~0.0040%, Alt: 0.010~0.040%, 0<B≤0.0004%, 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, 0<Sn≤0.0050%, 0<Sb≤0.0025%, 0<Bi≤0.0030%, 0<Pb≤0.0050%, 0<As≤0.0030%, the balance is Fe and unavoidable impurities, and meets the following requirements:

[0018] Nb+Ti≤0.07%;

[0019] Ca / S≥2;

[0020] Cu+Ni+Cr+Mo≤0.80%;

[0021] 0<N+O+H≤0.0070%;

[0022] 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0; and

[0023] The hydrogen-induced cracking sensitivity coefficient Ks≤0.23, Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19, where the element symbol is substituted into the value before the mass percentage sign of the corresponding element content.

[0024] The chemical composition design principle of the pipeline steel described in the present invention is as follows:

[0025] Carbon (C) is the most economical strengthening element in steel, increasing its strength through interstitial solid solution strengthening. Increasing the carbon content significantly improves the steel's hardenability, reduces the addition of other precious alloying elements, and lowers production costs. However, increased C content negatively impacts the steel's ductility, toughness, weldability, and corrosion resistance, particularly by exacerbating core segregation and significantly negatively impacting acid resistance. Therefore, the present invention adopts an ultra-low C design, with its mass percentage controlled between 0.015% and 0.040%, preferably between 0.017% and 0.032%.

[0026] Si: Si is a solid solution strengthening element and also acts as a deoxidizing element in steel. However, excessive Si content can adversely affect surface quality and weldability. If the Si content exceeds 0.30%, toughness may decrease. Therefore, in the present invention, the Si content is controlled within a range of 0.10-0.30% by weight, preferably 0.18-0.27%.

[0027] Mn: Mn increases steel strength through solid solution strengthening and is the most important and economical strengthening element in steel for compensating for the strength loss caused by reduced carbon content. Mn helps produce fine phase transformation products and also assists in controlling oxygen and sulfur during the steelmaking process. If the Mn content is too low, it will be difficult to achieve the target strength level, while too high a Mn content will also aggravate center segregation. Therefore, in the present invention, the Mn content is controlled within a range of 1.18% to 1.45% by weight, preferably 1.25% to 1.43%.

[0028] Cr: Cr has a certain solid solution strengthening effect and can effectively improve the hardenability of steel. A Cr content of 0.15% or above can effectively improve the corrosion resistance of steel, forming a relatively dense protective layer on the steel surface, which protects the substrate. However, when the Cr content in the steel is too high, it will be detrimental to the quality of the weld and easily form gray spot defects. Therefore, in the present invention, the Cr content by weight is controlled to 0.15-0.30%, preferably 0.20-0.25%.

[0029] Nb: Nb is an important element in low-carbon microalloyed steel. During hot rolling, the solid-soluted Nb precipitates through strain induction to form Nb carbonitrides, which pin the grain boundaries and inhibit the growth of deformed austenite. After controlled rolling and controlled cooling, the deformed austenite is transformed into fine products with a high dislocation density. After the steel strip is coiled, the solid-solution Nb will disperse and precipitate in the matrix as second-phase particles NbC, playing a role in precipitation strengthening. Excessive Nb makes the slab prone to cracking, thereby affecting the surface quality and also deteriorating the welding performance. Therefore, in the present invention, the mass percentage of Nb is controlled to 0.025-0.055%, preferably 0.035-0.045%.

[0030] Ti: Ti is an effective deoxidizer and degasser, and an effective element for fixing nitrogen and carbon. Undissolved Ti carbonitrides prevent austenite grain growth when steel is heated. TiN and TiC precipitated during rough rolling in the high-temperature austenite region effectively inhibit austenite grain growth, thereby refining the grains. Ti also increases the solid solubility of Nb, reducing the microcrack sensitivity of Nb-containing steels. Ti is typically added in combination with Nb. Furthermore, precipitation during welding can inhibit high-temperature grain growth, thereby improving weldability. Therefore, in the present invention, the Ti content by mass is controlled to 0.005 to 0.020%, preferably 0.007 to 0.018%.

[0031] Nb+Ti: Adding too much Nb and Ti can easily lead to the aggregation of coarse Nb / Ti particles in the steel plate, and easily form long strip-shaped segregations distributed along the rolling direction in the core. The long strip-shaped segregations are hard phases and have higher hardness than the matrix, which increases the hydrogen-induced cracking sensitivity of the structure, causing hydrogen to be pinned around them and no longer diffuse. Therefore, they are likely to become hydrogen aggregation sites and the starting point of hydrogen-induced cracking, significantly increasing the HIC cracking tendency of the core and deteriorating the acid resistance of the steel. Therefore, in the present invention, the mass percentage content of Nb and Ti elements is controlled to satisfy Nb+Ti≤0.07%, preferably 0.033~0.069%, and more preferably 0.033~0.065%.

[0032] In addition, TiN and NbC are important secondary phases in steel, significantly impacting strength and toughness. When (Nb+27.5*Ti) / (5C+27*N)>5, the secondary phase particles are coarse and unevenly distributed, easily concentrating at the core segregation site, causing stress concentration and brittle fracture. This leads to a sharp decrease in the steel's low-temperature toughness, significantly reducing the toughness of the steel and the hot zone of the weld, and also becoming a hard phase that causes hydrogen-induced cracking. When (Nb+27.5*Ti) / (5C+27*N)<0.7, the number of precipitated phases is too small and their size is too small, making them easily dissolved into the steel again. The grain boundary pinning effect is weak, which is also detrimental to toughness. Therefore, in the present invention, the mass percentage content of Nb and Ti elements is controlled to satisfy 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0, preferably 0.81≤(Nb+27.5*Ti) / (5C+27*N)≤3.69, and more preferably 1.13≤(Nb+27.5*Ti) / (5C+27*N)≤2.70.

[0033] Mo: Mo is a strong hardenability element. It significantly delays ferrite transformation, inhibits the formation of ferrite and pearlite, effectively promotes bainite transformation, and strengthens the matrix, resulting in a finer structure. A certain amount of Mo significantly improves the structural properties of steel. However, excessive Mo content reduces the steel's plasticity and increases its cost. Therefore, in the present invention, the Mo content is controlled to a low level of 0.02-0.09% by weight.

[0034] Cu and Ni: Cu and Ni can increase the strength of steel through solid solution strengthening and improve its atmospheric corrosion resistance. However, excessive Cu and Ni content can lead to hot cracking in the steel plate. Ni, on the other hand, can refine the grain size and improve the hot brittleness easily caused by Cu in steel, significantly contributing to the improvement of low-temperature toughness. Therefore, in the present invention, the mass percentage of Cu is controlled within a range of 0.05 to 0.35%, preferably 0.15 to 0.33%, and the mass percentage of Ni is controlled within a range of 0.05 to 0.30%, preferably 0.10 to 0.22%.

[0035] Cu+Ni+Cr+Mo: The addition of Cu, Ni, Cr, and Mo increases production costs. Excessive Cu, Ni, Cr, and Mo content can increase the steel's tendency to hot brittleness. It can also promote the formation of MA particles and increase the proportion of hard phases in the steel, acting as hydrogen traps. This can lead to hydrogen embrittlement and brittle phases, reducing HIC resistance and low-temperature toughness, impacting the steel's performance. To ensure the steel's overall strength, toughness, and other properties while reducing production costs, the mass percentages of Cu, Ni, Cr, and Mo in the present invention are controlled to satisfy Cu+Ni+Cr+Mo≤0.80%, preferably 0.34≤Cu+Ni+Cr+Mo≤0.79%.

[0036] S and P: S and P are major impurity elements in steel. P can easily cause cold brittleness, while S can easily cause hot brittleness, leading to unstable steel properties. In particular, as the S content increases, MnS inclusions increase, significantly increasing the steel's susceptibility to hydrogen-induced cracking and reducing its acid resistance. Therefore, the phosphorus and sulfur contents in steel should be minimized. Therefore, in the present invention, the mass percentages of S and P are controlled to S ≤ 0.0010% and P ≤ 0.0080%.

[0037] Ca: Ca treatment can control the morphology of sulfides, improve the anisotropy of the steel plate, and enhance low-temperature toughness. Furthermore, when the S content is too high, maintaining an appropriate Ca / S ratio ensures the steel's HIC resistance. However, excessively high Ca content can easily lead to an increase in inclusions such as CaO. In the present invention, the Ca content is controlled within a range of 0.0010 to 0.0040% by mass, preferably 0.0021 to 0.0035%, while the Ca / S ratio is controlled to be ≥ 2. Preferably, the Ca / S ratio is 3 to 13, and more preferably 4.1 to 8.7.

[0038] Al: Al is used to deoxidize steel. Appropriate amounts of Al also help refine grains and improve strength and toughness. However, if the total aluminum (Alt) content exceeds 0.04%, coarse precipitates may form, reducing the steel's acid resistance and low-temperature toughness. Therefore, the present invention controls the Al content to 0.010-0.040% by weight.

[0039] N, O, H: For low-temperature toughness, O, N, and H elements may all cause brittle fracture, especially at low temperatures. The presence of O and N will form oxygen compounds and nitrogen compounds, which may aggregate at grain boundaries, leading to intergranular brittleness. At the same time, H may also penetrate into the grains, causing the lattice to deform and form holes, thereby leading to brittle fracture. Therefore, the low-temperature toughness of steel is related to the content of O, N, and H. Excessive or inappropriate content may lead to brittle fracture. For acid resistance, the effects of O, N, and H are also quite significant. In an acidic environment, O will react with the iron in the steel to oxidize it, causing the steel to lose a certain degree of corrosion resistance. N may form nitrogen compounds with the chromium in the steel, resulting in a decrease in the corrosion resistance of the steel. H will promote stress corrosion cracking of the steel and will also have a certain impact on the acid resistance of the steel. Therefore, in the present invention, the mass percentage contents of N, O, and H are controlled to be 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, and 0<N+O+H≤0.0070%, preferably 0.0024%≤N+O+H≤0.0065%.

[0040] B: The primary function of B in steel is to increase its hardenability and strength, thereby conserving other rare and precious metals. However, its addition can significantly negatively impact the steel's low-temperature toughness. Therefore, in the present invention, the B content is controlled to 0 < B ≤ 0.0004% by weight.

[0041] Sn, Sb, Bi, Pb, As: When the content of Sn, Sb, Bi, Pb, and As in steel exceeds a certain limit, it will reduce the strength of the steel and easily cause severe segregation, forming segregation at grain boundaries. This causes brittle cracks to initiate from the grain boundaries and extend along the grain boundaries until complete fracture, significantly reducing the low-temperature toughness of the steel. In addition, excessive Sn content in steel will significantly reduce the steel's resistance to HIC hydrogen-induced cracking. This is because Sn promotes the diffusion and aggregation of hydrogen, thereby accelerating the occurrence of hydrogen-induced cracking. Pb has a similar effect. Therefore, in the present invention, the mass percentage content of Sn, Sb, Bi, Pb, and As elements is controlled to 0 < Sn ≤ 0.0050%, 0 < Sb ≤ 0.0025%, 0 < Bi ≤ 0.0030%, 0 < Pb ≤ 0.0050%, and 0 < As ≤ 0.0030%.

[0042] Hydrogen-induced cracking sensitivity coefficient (Ks): HIC resistance is affected by various elements in the steel, including C and S, which have significant segregation, Nb and Ti, which participate in the precipitation of second-phase particles, and inclusions such as N and B. Based on the segregation potential of each element and research on the relationship between each element and HIC resistance, the present invention proposes a hydrogen-induced cracking sensitivity coefficient (Ks): Ks = 8.0C + 24S - Alt - 2.3Nb + 2.8Ti + 117B + 0.34Ni - 0.30Mo + 24.5N - 0.19, where the symbol for each element represents the mass percentage of the element before the element's content. The correlation between the HIC crack length ratio (CLR) and Ks is shown in Figure 1. As Ks increases, the steel's hydrogen-induced cracking sensitivity increases, the probability of HIC cracks increases, and the crack length increases. When Ks ≤ 0.23, no steel undergoing HIC testing cracks. When Ks > 0.23, the CLR begins to increase significantly. The API 5L SPEC standard requires CLR to be ≤15%, while some other engineering projects require CLR ≤10% or even more stringent. Based on the relationship between CLR and Ks, in order to ensure the acid resistance of steel, the present invention controls Ks to be ≤0.23.

[0043] In general, the present invention adopts the above-designed composition system, comprehensively considering the effects of easily segregated elements such as C and S, microalloying elements such as Cu, Ni, Cr, Mo, Nb, and Ti, as well as elements such as N, O, H, B, Sn, Sb, Bi, Pb, and As in steel on the acid resistance and low-temperature toughness of the steel. Based on the degree of segregation of each element and its relationship with HIC resistance, the present invention proposes control requirements for the hydrogen-induced cracking sensitivity coefficient Ks to improve the acid resistance of the steel. On this basis, the focus is on considering the effects of second-phase particles such as TiN and NbC on the acid resistance and low-temperature toughness of steel, controlling Nb+Ti≤0.07%, 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0, and at the same time, by controlling Cu+Ni+Cr+Mo≤0.80%, reducing the inhibitory effect on massive ferrite, so that a small amount of massive ferrite is contained in the steel structure (the massive ferrite content is 3-10 volume%), thereby suppressing the generation of cracks, so that the pipeline steel has good low-temperature toughness on the basis of high strength and acid resistance.

[0044] Preferably, the microstructure of the pipeline steel is granular bainite + massive ferrite.

[0045] Preferably, the content of the massive ferrite is 3 to 10 volume %, more preferably 4.6 to 8.5 volume %.

[0046] Preferably, the grain size of the massive ferrite is ≥10, more preferably 10.5-11.

[0047] Preferably, the content of the granular bainite is 90 to 97 volume %, more preferably 91.5 to 95.4 volume %.

[0048] Preferably, the granular bainite is uniformly refined granular bainite.

[0049] Preferably, there is no obvious segregation band or hard phase structure in the microstructure of the pipeline steel.

[0050] Preferably, the R of the pipeline steel t0.5 The yield strength is 430 MPa or more, preferably 450-570 MPa, more preferably 510-570 MPa; the tensile strength is 520 MPa or more, preferably 540-690 MPa, more preferably 615-690 MPa; the yield strength ratio is ≤0.92, preferably ≤0.87, more preferably 0.81-0.87; the elongation A50 is ≥30%, preferably 30%-47%, more preferably 40%-47%; the hardness HV10 is ≤220, preferably 195-220; the Charpy impact energy at -49°C is ≥400 J, preferably ≥430 J, more preferably 430 J-500 J; the DWTT SA% at -35°C is ≥90%, preferably 90-98%.

[0051] Preferably, the pipeline steel meets at least one of the following properties: no cracking in a HIC resistance test conducted according to the NACE TM0284-2016 standard "Test Method for Evaluation of Hydrogen Induced Cracking Resistance of Steel for Pipelines and Pressure Vessels"; no cracking in a SSC resistance test conducted according to the NACE TM0177-2016 standard "Test Method for Resistance of Metallic Materials to Sulfide Stress Corrosion Cracking in Pressurized Sulfur Environments".

[0052] On the other hand, the present invention also provides a method for manufacturing the above-mentioned pipeline steel, which comprises the following steps:

[0053] 1) Hot metal pretreatment

[0054] Control the S content of molten iron before pretreatment to ≤0.050% and the S content of molten iron entering the furnace to ≤0.0020%;

[0055] 2) Smelting

[0056] The molten iron in step 1) is smelted to control the carbon content of the tapped steel to be ≤0.035% and the sulfur content of the tapped steel to be ≤0.0080%;

[0057] 3) Refining

[0058] The steel tapped in step 2) is subjected to LF refining and RH refining. During LF refining, the stirring time after all alloy additions is ≥3 minutes. During RH refining, the amount of calcium wire fed is 200-350m, and the argon blowing time after the calcium wire feeding is completed is ≥5 minutes. The target Ca content is 10-40ppm, and the Ca / S ratio is ≥2.

[0059] 4) Continuous casting

[0060] The cooling time of continuous casting should be ≥3min, and the superheat should be controlled at 10~40℃;

[0061] 5) Reheating

[0062] The heating temperature is 1100-1200°C and the holding time is 200-350 minutes;

[0063] 6) Controlled rolling

[0064] The final rolling temperature is 750-900℃;

[0065] 7) Cooling after rolling

[0066] The start cooling temperature is 700-800°C, the stop cooling temperature is 300-550°C, the cooling rate is ≤30°C / s, and the cooling rate is preferably 10-30°C / s.

[0067] Preferably, in step 2), the bottom blowing flow rate of the converter bottom calming stage is ≥900Nm 3 / h, control the blowing of free oxygen to ≤900ppm.

[0068] Composition design and internal quality control are crucial for improving the acid resistance and toughness of steel. Increased S content in the process and finished product increases the precipitation and aggregation of inclusions such as MnS in the steel. MnS inclusions generated in the core are elongated during rolling and easily form gaps after cooling. MnS is one of the most vulnerable locations for HIC formation, significantly reducing low-temperature toughness and HIC resistance. To ensure low S content in the finished product and process, in step 1) of the present method, the S content of the pre-treatment molten iron is controlled to ≤0.050%, and the S content of the incoming molten iron is controlled to ≤0.0020%. After desulfurization, the molten iron ladle is cleanly skimmed.

[0069] In step 2) of the method of the present invention, during the smelting process, the bottom blowing flow rate of the converter bottom calming stage is ≥900Nm 3 / h, thereby enhancing molten pool agitation, making the composition and temperature more uniform, and improving molten steel quality; controlling the free oxygen content to ≤900ppm when blowing is stopped ensures a low oxygen content in the molten steel during the shutdown period, reducing defects such as pores and inclusions during subsequent solidification, and improving slab quality. Increased carbon content exacerbates segregation in the steel, further aggregating inclusions such as MnS in the core and significantly increasing the tendency to hydrogen-induced cracking. Therefore, the carbon and sulfur content of the ladle steel must be strictly controlled. In the present invention, the carbon content of the ladle steel (i.e., the tapped steel) is controlled to C≤0.035% and S≤0.0080%.

[0070] In step 3) of the present method, during LF refining, aluminum slag is used for slag deoxidation, depending on the pre-desulfurization and pre-deoxidation conditions after the converter. Desulfurization is achieved through bottom stirring with argon, with a stirring time of ≥3 minutes. During RH refining, the target feed rate of calcium silicon wire is 200-350 m3. After the RH refining is completed, a low-flow ladle bottom argon blowing is performed for at least 5 minutes, with a target Ca content of 10-40 ppm and a Ca / S ratio of ≥2 to modify inclusions such as MnS and Al2O3 in the steel. By increasing the effective calcium content in the steel, high-melting-point CaS is formed earlier during the solidification process, suppressing the total amount of MnS generated during this process and converting it entirely or partially into CaS. Furthermore, large Al2O3 inclusions are converted into low-melting-point composite inclusions, promoting their buoyancy and purifying the molten steel. Both LF and RH refineries utilize low-carbon alloys, with a C content of 0.015-0.040%.

[0071] In step 4) of the present method, during the continuous casting process, the continuous casting quieting time is controlled to ≥ 3 minutes to reduce defects such as shrinkage cavities and porosity in the slab and improve slab quality. A high superheat can increase segregation, while a too low superheat prevents inclusions from fully floating, both of which negatively impact acid resistance and low-temperature toughness. Therefore, in the present invention, the superheat is controlled between 10 and 40°C to control center segregation, with a low-magnification rating of ≤ 2 according to the Mannesmann standard. Preferably, a full-process protection and constant-rate casting method is employed to ensure good internal quality and segregation levels in the slab.

[0072] In step 5) of the method of the present invention, the slab is heated to a temperature of 1100-1200°C and held at this temperature for 200-350 minutes. Sufficient heating temperature and time ensure sufficient solid solution of the alloy and good uniformity, thereby achieving a uniform microstructure and plate shape.

[0073] In step 6) of the method of the present invention, the finishing rolling temperature is controlled between 750°C and 900°C. If the finishing rolling temperature is too high, the austenite grains will coarsen significantly, which will adversely affect the microstructure and properties of the final steel plate. If the finishing rolling temperature is too low, mixed grains will easily form in the steel, adversely affecting the toughness, plasticity, and processability of the steel. Preferably, the rolling process is performed using the TMCP process.

[0074] In step 7) of the present method, a DQ+ACC cooling mode is adopted after rolling. By limiting the addition of Mo and Nb elements to relatively low levels, their inhibitory effect on bulk ferrite is appropriately reduced. This allows the steel to undergo phase transformation under conditions of a start cooling temperature of 700-800°C, a stop cooling temperature of 300-550°C, and a cooling rate of ≤30°C / s (preferably 10-30°C / s). At this time, a small amount of ferrite first precipitates along the original austenite grain boundaries and has sufficient phase transformation time to transform into relatively regular bulk ferrite, thereby obtaining a small amount of bulk ferrite (≤10 volume %) without affecting the phase transformation of bainite. As a tough phase, bulk ferrite can increase the interface area and cause significant plastic deformation at the crack tip, alleviating the triaxial tensile stress at the crack tip and hindering crack propagation. At the same time, the presence of granular bainite inhibits the coordinated deformation of ferrite and increases deformation resistance. Most cracks need to frequently cross the interface between bainite and polygonal ferrite during their propagation process. When a small proportion of blocky ferrite is present in the structure, crack development is effectively suppressed, resulting in better low-temperature DWTT toughness. Preferably, the content of blocky ferrite is 3-10% by volume. When it is less than 3% by volume, the structure's ability to coordinate deformation is insufficient, the overall toughness and ductility are poor, and stress concentration is easily generated, leading to rapid crack propagation. When it is greater than 10% by volume, the structure's deformation resistance is low, the overall strength and toughness are insufficient, and the ability to resist crack development is poor.

[0075] In addition, a cooling rate of 10 to 30°C / s is conducive to rapid and uniform transformation of the structure, reducing core structure segregation and inclusion aggregation, and avoiding local hard point defects on the surface, obtaining uniformly refined structure grains in the thickness direction and throughout the plate, ensuring uniform and stable overall performance of the steel plate, and ensuring that the steel plate has excellent acid resistance while having good low-temperature toughness.

[0076] Preferably, the method satisfies at least one of the following conditions:

[0077] In step 2), the C content of the tapped steel is 0.018% to 0.035%, and / or the S content of the tapped steel is 0.0036% to 0.0080%;

[0078] In step 3), the stirring time after all the alloys are added is 3 to 12 minutes, the argon blowing time after the calcium line is finished is 5 to 14 minutes, and / or the Ca / S ratio is 3 to 13;

[0079] In step 4), the continuous casting calming time is 3 to 12 minutes;

[0080] In step 7), the cooling rate is 15-30°C / s.

[0081] Compared with the prior art, the present invention has the following advantages:

[0082] In terms of composition design, the present invention adopts an ultra-low C, ultra-low S and low inclusion element scheme. Based on the degree of segregation of each element and its relationship with HIC resistance, it proposes control requirements for the hydrogen-induced cracking sensitivity coefficient Ks (i.e., Ks≤0.23, Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19) to improve the acid resistance of the steel. On this basis, the Nb+Ti content is controlled to ≤ 0.07%, and 0.7 ≤ (Nb+27.5*Ti) / (5C+27*N) ≤ 5.0, with a focus on the effects of second-phase particles such as TiN and NbC on the acid resistance and low-temperature toughness of the steel. At the same time, by controlling the Cu+Ni+Cr+Mo content to ≤ 0.80%, the inhibitory effect on massive ferrite is reduced, allowing the steel structure to contain a small amount of massive ferrite (3-10%), suppressing the generation of cracks. This ensures that the pipeline steel not only has high strength and acid resistance but also good low-temperature toughness. Conventional acid-resistant pipeline steels, on the other hand, contain high levels of C, Mo, and Nb, and do not consider the control of inclusion elements (such as N and B).

[0083] In terms of manufacturing method, the present invention adopts DQ+ACC cooling mode after rolling, controls the start cooling temperature to 700-800°C, the stop cooling temperature to 300-550°C, and the cooling rate ≤30°C / s (preferably 10-30°C / s). At the same time, since the inhibitory effect of elements such as Mo on block ferrite is reduced, a microstructure of uniformly refined granular bainite + a small amount of block ferrite can be formed in this mode, wherein the block ferrite content is 3-10%, the grain size is ≥10 levels, and there is no obvious segregation band or hard phase structure. Block ferrite, as a tough phase, can increase the interface area and cause greater plastic deformation at the crack tip, thereby alleviating the triaxial tensile stress at the crack tip and hindering crack propagation, thereby making the acid-resistant pipeline steel have good low-temperature toughness.

[0084] The R of the pipeline steel of the present invention t0.5 The pipeline steel has a yield strength of 430-570 MPa, a tensile strength of 520-690 MPa, a yield strength ratio of ≤0.92, an elongation A50 of ≥30%, a hardness HV10 of ≤220, a -49°C Charpy impact energy of ≥400 J, and a -35°C DWTT SA% of ≥90%. It exhibits excellent acid resistance (no cracking in HIC and SSCC tests (90% SYMS) conducted in NACE Standard A solution), fully meeting the API 5L assessment standard. Compared to existing technologies, the pipeline steel of the present invention exhibits improved low-temperature toughness, with a -49°C Charpy impact energy of ≥400 J and a -35°C DWTT SA% of ≥90%. Among conventional acid-resistant pipeline steels, the highest low-temperature toughness is achieved with a -40°C Charpy impact energy of ≥300 J and a -30°C DWTT SA% of ≥85%. BRIEF DESCRIPTION OF THE DRAWINGS

[0085] Figure 1 shows the scatter plot of the HIC crack length rate (CLR) and the coefficient Ks. As shown in the figure, when Ks ≤ 0.23, the CLR is 0, meaning no cracking occurs.

[0086] FIG2 is a microstructure photograph of the pipeline steel according to Example 1 of the present invention. DETAILED DESCRIPTION

[0087] The pipeline steel and the manufacturing method thereof of the present invention will be further described below with reference to the embodiments and the accompanying drawings. This description does not limit the technical solution of the present invention.

[0088] The pipeline steels of Examples 1-10 of the present invention and the comparative steels of Comparative Examples 1-2 were both prepared by the following steps:

[0089] 1) Hot metal pretreatment

[0090] Control the S content of molten iron before pretreatment to ≤0.050% and the S content of molten iron entering the furnace to ≤0.0020%;

[0091] 2) Smelting

[0092] For smelting molten iron, the bottom blowing flow rate of the converter bottom calming stage is ≥900Nm 3 / h, control the free oxygen to ≤900ppm, control the C content of the tapped steel to ≤0.035%, and the S content of the tapped steel to ≤0.0080%;

[0093] 3) Refining

[0094] The steel tapped in step 2) is subjected to LF refining and RH refining. During LF refining, the stirring time after all alloy additions is ≥3 minutes. During RH refining, the amount of calcium wire fed is 200-350m, and the argon blowing time after the calcium wire feeding is completed is ≥5 minutes. The target Ca content is 10-40ppm, and the Ca / S ratio is ≥2.

[0095] 4) Continuous casting

[0096] The cooling time of continuous casting should be ≥3min, and the superheat should be controlled at 10~40℃;

[0097] 5) Reheating

[0098] The heating temperature is 1100-1200°C and the holding time is 200-350 minutes;

[0099] 6) Controlled rolling

[0100] Finish rolling temperature: 750~900℃;

[0101] 7) Cooling after rolling

[0102] After rolling, the DQ+ACC cooling mode is adopted, with the start cooling temperature of 700-800℃, the stop cooling temperature of 300-550℃, and the cooling rate of 10-30℃ / s.

[0103] The chemical element compositions of the pipeline steels of Examples 1-10 of the present invention and the comparative steels of Comparative Examples 1-2, produced through the above steps, are shown in Tables 1 and 2, with the remainder comprising Fe and unavoidable impurities. The process parameters for the manufacturing methods of the pipeline steels of Examples 1-10 of the present invention and the comparative steels of Comparative Examples 1-2 are shown in Table 3. The chemical element compositions and related process designs of the pipeline steels of Examples 1-10 meet the design requirements of the present invention. The comparative steel of Comparative Examples 1-2 was produced using a method substantially identical to that of Examples 1-10 of the present invention, except that at least one of its chemical composition and / or process parameters did not meet the requirements of the present invention.

[0104] The pipeline steels of Examples 1-10 and the comparative steels of Comparative Examples 1-2 were sampled and subjected to various performance tests. The obtained mechanical properties and microstructure test results are shown in Table 4, and the HIC resistance and SSC performance test results are shown in Table 5.

[0105] The relevant performance test methods are as follows:

[0106] (1) Tensile, hardness, and Charpy impact tests: Tensile, hardness, and Charpy impact tests were performed according to ASTM A370 to obtain the R values ​​of the pipeline steels of each embodiment and the comparative steels of each comparative example at room temperature. t0.5 Yield strength, tensile strength, yield strength ratio, elongation A50, HV10 Vickers hardness, and -49°C Charpy impact energy value. Yield strength ratio = yield strength / tensile strength.

[0107] (2) DWTT drop weight tear test: The -35°C DWTT SA% of the pipeline steels of each embodiment and the comparative steels of each comparative example were obtained according to API RP 5L3.

[0108] (3) Microstructure observation: The microstructure morphology of Example 1 was observed by an optical microscope (manufacturer: ZEISS, model: Axio Imager.M2m), and the volume percentage content and grain size of the bulk ferrite were measured by metallographic microscopy.

[0109] (4) Hydrogen-Induced Cracking (HIC) Resistance Test: The NACE TM0284-2016 standard, "Test Method for Evaluation of Hydrogen-Induced Cracking Resistance of Steel for Pipelines and Pressure Vessels," was used to measure the hydrogen-induced cracking (HIC) resistance of the steel in a wet hydrogen sulfide (H2S) environment. The steel was immersed in the standard sample and solution A for 96 hours. The measurement results are shown in Table 5. None of the samples obtained from Examples 1 to 10 of the present invention showed any cracks.

[0110] (5) Hydrogen sulfide stress corrosion cracking (SSC) resistance test: According to the NACE TM 0177-2016 standard "Test method for resistance of metallic materials to sulfide stress corrosion cracking in pressurized sulfur environments", the hydrogen sulfide stress corrosion cracking (SSC) resistance was measured using a four-point bending method with a stress loading of 90% SYMS. After 720 hours of immersion in solution A, the measurement results are shown in Table 5. None of the samples obtained in the embodiments of the present invention broke.

[0111] Table 5

[0112] As shown in Table 4, the pipeline steel obtained by the present invention has an Rt0.5 yield strength of 430-570 MPa, a tensile strength of 520-690 MPa, a yield ratio ≤0.92, an elongation A50 ≥30%, a hardness HV10 ≤220, and good low-temperature toughness, with a Charpy impact energy ≥400 J at -49°C and a DWTT SA% ≥90% at -35°C. The massive ferrite content is 3-10% by volume, and the grain size is ≥10. FIG2 also shows a microstructure photograph of the pipeline steel obtained by Example 1 of the present invention. As shown in FIG2, the steel plate obtained using the chemical composition design and manufacturing method of the present invention has a microstructure of uniformly refined granular bainite and a small amount of massive ferrite. The massive ferrite content is 3-10%, the grain size is ≥10, and there is no obvious segregation band or hard phase structure.

[0113] From the test results in Table 5, it can be seen that no cracking occurred in the pipeline steels of Examples 1-10 in the hydrogen induced cracking (HIC) resistance test and the hydrogen sulfide stress corrosion cracking (SSC) resistance test.

[0114] In the comparative steel of Comparative Example 1, the C content is 0.042 (exceeding the range C of the present invention: 0.015~0.040), and Ks reaches 0.311 (Ks≤0.23 in the present invention), (Nb+27.5*Ti) / (5C+27*N) is 0.61 (0.7~5.0 in the present invention), and Cu+Ni+Cr+Mo reaches 0.93% (Cu+Ni+Cr+Mo≤0.80% in the present invention). The comparative steel of Comparative Example 1 has large cracks in the HIC test, two of the three specimens are cracked, and the crack length rate CLR exceeds 20%. Cracks also appear in the SSC test, and the DWTT and impact properties are worse than those in the embodiment.

[0115] In the comparative steel of Comparative Example 2, Ks reaches 0.289 (Ks of the present invention is ≤0.23), the length of the calcium feeding wire is relatively small, and the Ca / S ratio is only 1.0 (Ca / S of the present invention is ≥2). In addition, the B content reaches 0.0005% (B content of the present invention is ≤0.0004%), and O+N+H reaches 0.0072% (O+N+H of the present invention is ≤0.0070%). The comparative steel of Comparative Example 2 has poor HIC resistance and toughness. The crack length rate CLR of one specimen reaches 18.56%, cracks also appear in the SSC test, the DWTT is only 70% / 75%, the average impact energy is 279 J, and one of the single values ​​is only 208 J.

[0116] In summary, the pipeline steel of the present invention can achieve the excellent effect of having high strength and acid resistance while also having good low-temperature toughness through reasonable component matching and process design.

[0117] It should be noted that the combination of the various technical features in this case is not limited to the combination described in the claims of this case or the combination described in the specific embodiments. All technical features recorded in this case can be freely combined or combined in any way unless there is a contradiction between them.

[0118] It should also be noted that the embodiments listed above are merely specific embodiments of the present invention. Obviously, the present invention is not limited to the above embodiments, and similar changes or modifications made therefrom that can be directly derived from or easily conceived by those skilled in the art based on the disclosure of the present invention are intended to fall within the scope of protection of the present invention.

Claims

1. A pipeline steel, which contains, in addition to Fe and unavoidable impurities, the following chemical elements in percentage by mass: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0<P≤0.008%, 0<S≤0.0010%, Cu: 0.05-0.35%, Ni: 0.05-0.30%, Cr: 0.15-0.30%, Mo: 0.02-0.09%, Nb: 0.025-0 .055%, Ti: 0.005-0.020%, Ca: 0.0010-0.0040%, Alt: 0.010-0.040%, 0<B≤0.0004%, 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, 0<Sn≤0.0050%, 0<Sb≤0.0025%, 0<Bi≤0.0030%, 0<Pb≤0.0050%, 0<As≤0.0030%, and meet the following requirements: Nb+Ti≤0.07%; Ca / S≥2; Cu+Ni+Cr+Mo≤0.80%; 0<N+O+H≤0.0070%; 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0; The hydrogen-induced cracking sensitivity coefficient Ks≤0.23, Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19, where the element symbol is substituted into the value before the mass percentage sign of the corresponding element content.

2. The pipeline steel according to claim 1, wherein: The pipeline steel contains the following chemical elements by mass percentage: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0<P≤0.008%, 0<S≤0.0010%, Cu: 0.05-0.35%, Ni: 0.05-0.30%, Cr: 0.15-0.30%, Mo: 0.02-0.09%, Nb: 0.025-0.055%, Ti: 0.005-0. 0.20%, Ca: 0.0010~0.0040%, Alt: 0.010~0.040%, 0<B≤0.0004%, 0<O≤0.0026%, 0<N≤0.0050%, 0<H≤0.0002%, 0<Sn≤0.0050%, 0<Sb≤0.0025%, 0<Bi≤0.0030%, 0<Pb≤0.0050%, 0<As≤0.0030%, the balance is Fe and unavoidable impurities, and meets: Nb+Ti≤0.07%; Ca / S≥2; Cu+Ni+Cr+Mo≤0.80%; 0<N+O+H≤0.0070%; 0.7≤(Nb+27.5*Ti) / (5C+27*N)≤5.0; The hydrogen-induced cracking sensitivity coefficient Ks≤0.23, Ks=8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19, where the element symbol is substituted into the value before the mass percentage sign of the corresponding element content.

3. The pipeline steel according to claim 1 or 2, wherein: The microstructure of the pipeline steel is granular bainite+blocky ferrite.

4. The pipeline steel according to claim 3, wherein: The content of the massive ferrite is 3 to 10 volume %, preferably 4.6 to 8.5 volume %.

5. The pipeline steel according to claim 3 or 4, wherein: The grain size of the pipeline steel is ≥ grade 10, preferably 10.5-11.

6. The pipeline steel according to any one of claims 1 to 5, wherein: The R of the pipeline steel t0.5 The yield strength is above 430MPa, preferably 450-570MPa; the tensile strength is above 520MPa, preferably 540-690MPa; the yield strength ratio is ≤0.92, the elongation A50 is ≥30%, preferably ≥40%; the hardness HV10 is ≤220, the Charpy impact energy at -49°C is ≥400J, and the DWTT SA% at -35°C is ≥90%.

7. The pipeline steel according to any one of claims 1 to 6, wherein: The pipeline steel meets at least one of the following properties: no cracks in the HIC resistance test conducted according to the NACE TM0284-2016 standard "Test method for evaluation of hydrogen-induced cracking resistance of steel for pipelines and pressure vessels"; no cracks in the SSC resistance test conducted according to the NACE TM 0177-2016 standard "Test method for resistance of metallic materials to sulfide stress corrosion cracking in pressurized sulfur environments".

8. A method for manufacturing the pipeline steel according to any one of claims 1 to 7, comprising the following steps: 1) Hot metal pretreatment Control the S content of molten iron before pretreatment to ≤0.050% and the S content of molten iron entering the furnace to ≤0.0020%; 2) Smelting The molten iron in step 1) is smelted to control the C content of the tapped steel to be ≤0.035% and the S content of the tapped steel to be ≤0.0080%; 3) Refining The steel produced in step 2) is subjected to LF refining and RH refining. During LF refining, the stirring time after all alloys are added is ≥3 min. During RH refining, the amount of calcium wire fed is 200-350 m, the argon blowing time after the calcium wire feeding is completed is ≥5 min, the target Ca content is 10-40 ppm, and the Ca / S ratio is ≥2. 4) Continuous casting Continuous casting calming time ≥ 3min, superheat controlled at 10~40℃; 5) Reheating The heating temperature is 1100-1200°C and the holding time is 200-350 minutes; 6) Rolling The final rolling temperature is 750~900℃; 7) Cooling after rolling The start cooling temperature is 700-800℃, the stop cooling temperature is 300-550℃, and the cooling rate is ≤30℃ / s.

9. The manufacturing method according to claim 8, wherein: In step 2), the bottom blowing flow rate of the converter bottom calming stage is ≥ 900Nm 3 / h, control the blowing of free oxygen to ≤900ppm.

10. The manufacturing method according to claim 8, wherein: The method satisfies at least one of the following conditions: In step 2), the C content of the tapped steel is 0.018% to 0.035%, and / or the S content of the tapped steel is 0.0036% to 0.0080%; In step 3), the stirring time after all the alloys are added is 3 to 12 minutes, the argon blowing time after the calcium wire feeding is completed is 5 to 14 minutes, and / or the Ca / S ratio is 3 to 13; In step 4), the continuous casting calming time is 3 to 12 minutes; In step 7), the cooling rate is 15-30°C / s.

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