Patent Grant
8821653
This application claims priority to Italian Patent Application No. MI2011A000179, entitled “HEAVY WALL STEEL PIPES WITH EXCELLENT TOUGHNESS AT LOW TEMPERATURE AND SULFIDE STRESS CORROSION CRACKING RESISTANCE”, filed Feb. 7, 2011. This application is also related to U.S. patent application Ser. No. 13/367,332, entitled “HIGH STRENGTH STEEL PIPES WITH EXCELLENT TOUGHNESS AT LOW TEMPERATURE AND SULFIDE STRESS CORROSION CRACKING RESISTANCE”, filed Feb. 6, 2012. The entirety of each of these applications is hereby incorporated by reference and should be considered a part of this specification.
1. Field
Embodiments of the present disclosure relate generally to metal production and, in certain embodiments, relates to methods of producing metallic tubular bars having high toughness at low temperature while concurrently possessing sulfide stress corrosion cracking resistance. Certain embodiments relate to heavy wall seamless steel pipes for risers, line pipes and flow lines for use in the oil and gas industry, including pipes that are suitable for bending.
2. Description of the Related Art
Exploration of offshore oil and gas reserves in remote regions of the world is increasingly moving away from conditions where relatively traditional pipe solutions can be utilized and towards more demanding environments. These more demanding environments may incorporate a combination of very challenging factors, including for example, deep water locations, increased pressure and temperature wells, more corrosive products, and lower design temperatures. These conditions, when added to stringent weldability and toughness criteria already associated with pipe specifications for offshore oil and gas exploration applications, place ever increasing demands on the materials and supply capability.
These demands are evident in project developments involving aggressive composition and high operating pressure that require very thick wall, sour service carbon steels. For example, major seamless line pipe manufacturers are able to manufacture pipes of grades X65 and X70 according with American Petroleum Institute (API) 5L and International Organization for Standards (ISO) 3183 standards, with sulfide stress corrosion (SSC) and hydrogen induced cracking (HIC) resistance, when wall thickness (WT) is below 35 mm. However, the conflicting requirements of strength and toughness, combined with the need for sulfide stress corrosion (SSC) and hydrogen induced cracking (HIC) resistance (e.g., sour resistance) in thick wall pipes (e.g., WT greater than or equal to 35 mm) has proven difficult to achieve.
In the complex scenario of line pipe projects for applications such as sour service, deep and ultra-deep water, Arctic-like areas, etc., heavy wall bends have also become an important feature of pipes.
Embodiments of the disclosure are directed to steel pipes or tubes and methods of manufacturing the same. In some embodiments, heavy wall, seamless, quenched and tempered (Q&T) steel pipes are provided having a wall thickness (WT) greater than or equal to about 35 mm and a minimum yield strength of about 65 ksi and about 70 ksi, respectively, with excellent low temperature toughness and corrosion resistance (sour service, H2S environment). The steel pipes are suitable for use as line pipes and risers, amongst other applications. In some embodiments, the seamless pipes are also suitable to produce bends of the same grade by hot induction bending and off-line quenching and tempering treatment. In one embodiment, the steel pipe has an outside diameter (OD) between about 6″ (152 mm) and about 28″ (711 mm), and wall thickness (WT) greater than about 35 mm.
In one embodiment, the composition of a seamless, low-alloy steel pipe may comprise (by weight):
about 0.05% to about 0.16% C;
about 0.20% to about 0.90% Mn;
about 0.10% to about 0.50% Si;
about 1.20% to about 2.60% Cr;
about 0.05% to about 0.50% Ni;
about 0.80% to about 1.20% Mo;
about 0.80% W maximum;
about 0.03% Nb maximum;
about 0.02% Ti maximum;
about 0.005% to about 0.12% V;
about 0.008% to about 0.040% Al;
about 0.0030% to about 0.012% N;
about 0.3% Cu maximum;
about 0.01% S maximum;
about 0.02% P maximum;
about 0.001% to about 0.005% Ca;
about 0.0020% B maximum;
about 0.020% As maximum;
about 0.005% Sb maximum;
about 0.020% Sn maximum;
about 0.030% Zr maximum;
about 0.030% Ta maximum;
about 0.0050% Bi maximum;
about 0.0030% O maximum; and
about, 0.00030% H maximum;
with the balance of the composition comprising iron and inevitable impurities.
The steel pipes may be manufactured into different grades. In one embodiment, a 450 MPa (65 ksi) grade steel pipe may be provided with the following properties:
In another embodiment, a 485 MPa (70 ksi) grade steel pipe may be provided with the following properties:
The steel pipe may have a minimum impact energy of about 200 J/150 J (average/individual) and a minimum average shear area of about 80% for both longitudinal and transverse Charpy V-notch (CVN) tests performed on standard size specimens at about −70° C. according to standard ISO 148-1. Embodiments of the steel pipe may also have a ductile-to-brittle transition temperature, measured by drop weight test (DWT) according to ASTM 208, lower than about −70° C. In one embodiment, the steel pipe may have a maximum hardness of about 248 HV10.
Steel pipes manufactured according to embodiments of the present disclosure may exhibit resistance to both hydrogen induced cracking (HIC) and sulfide stress corrosion (SSC) cracking. In one embodiment, discussed in greater detail below, HIC testing performed according with Standard TM0284-2003 Item No. 21215, using NACE solution A and test duration 96 hours, provided the following HIC parameters for the steel pipes when averaged on three sections of three test specimens:
Crack Length Ratio, CLR<5%
Crack Thickness Ratio, CTR=1%
Crack Sensitivity Ratio, CSR=0.2%
In another embodiment, SSC testing of samples of the above described steel pipes was performed in accordance with NACE TM0177, using test solution A and a test duration of about 720 hours. Under these conditions, the steel pipe samples exhibited no apparent failure at about 90% of the specified minimum yield stress (SMYS).
Steel pipes manufactured according to certain embodiments of the disclosure have a microstructure exhibiting substantially no ferrite, substantially no upper bainite, and substantially no granular bainite. The steel pipes may further comprise tempered martensite with volume percentage greater than about 50%, greater than about 60%, preferably greater than about 90%, and most preferably greater than about 95%, as measured according to ASTM E562-08, with tempered lower bainite in a volume percentage less than about 40%, preferably less than about 10%, most preferably less than about 5%. Martensite and bainite, in some embodiments, may be formed at temperatures lower than about 450° C. and about 540° C., respectively, after re-heating at temperatures of about 900° C. to about 1060° C. for soaking times between about 300 sec to about 3600 sec, and quenching at cooling rates equal or greater than about 7° C./s. In further embodiments, the average prior austenite grain size of the steel pipes, measured according to ASTM Standard E112, is greater than about 15 μm (lineal intercept) and smaller than about 100 μm. The average size of regions separated by high angle boundaries (i.e. packet size), in one embodiment, may be smaller than about 6 μm (preferably smaller than about 4 μm, most preferably smaller than about 3 μm), measured as average lineal intercept on images taken by Scanning Electron Microscopy (SEM) using the Electron Back Scattered Diffraction (EBSD) signal, and considering high-angle boundaries those with misorientation>45°. The microstructure may also include presence of one or more precipitates. For example, the microstructure may include precipitates of a first type given by any of MX and M2X, where M is selected from V, Mo, Nb, and Cr and X is selected from C and N. The size (e.g., average diameter) of the first type of precipitates may be less than about 40 nm. In further embodiments, coarse precipitates of a second type, given by any of M3C, M6C, M23C6, may also be present with the fine precipitates of the first type. The average diameter of the second type of precipitates may be between about 80 nm to about 400 nm (precipitates were examined by Transmission Electron Microscopy (TEM) using extraction replica method).
In one embodiment, a steel pipe is provided. The steel pipe comprises a steel composition comprising:
In another embodiment, a method of making a steel pipe is provided. The method comprises providing a steel having a carbon steel composition. The method further comprises forming the steel into a tube having a wall thickness greater than or equal to about 35 mm. The method additionally comprises heating the formed steel tube in a first heating operation to a temperature within the range between about 900° C. to about 1060° C. The method also comprises quenching the formed steel tube at a rate greater than or equal to about 7° C./sec, where the microstructure of the quenched steel is greater than or equal to about 50% martensite and less than or equal to about 50% lower bainite and has an average prior austenite grain size greater than about 15 μm. The method additionally comprises tempering the quenched steel tube at a temperature within the range between about 680° C. to about 760° C., where the steel tube, after tempering, has a yield strength greater than about 450 MPa (65 ksi) and a Charpy V-notch energy greater than or equal to about 150 J/cm2.
Other features and advantages of the disclosed embodiments will be apparent from the following description taken in connection with the accompanying drawings.
Embodiments of the present disclosure provide steel compositions and tubular bars (e.g., pipes) formed using the steel compositions. Respective methods of manufacture are also disclosed. The tubular bars may be employed, for example, as line pipes and risers for use in the oil and gas industry. In certain embodiments, the tubular bars may possess wall thicknesses greater than or equal to about 35 mm and a microstructure of martensite and lower bainite without substantial ferrite, upper bainite, or granular bainite. So formed, the tubular bars may possess a minimum yield strength of about 65 ksi and about 70 ksi. In further embodiments, the tubular bars may possess good toughness at low temperatures, as well as resistance to sulfide stress corrosion cracking (SSC) and hydrogen induced cracking (HIC). These properties enable use of the tubular bars in sour service environments. It may be understood, however, that tubular bars comprise one example of articles of manufacture which may be formed from embodiments of the present disclosure and should in no way be construed to limit the applicability of the disclosed embodiments.
The term “bar” as used herein is a broad term and includes its ordinary dictionary meaning and also refers to generally hollow, elongate members which may be straight or have bends or curves and may be formed to a predetermined shape, and any additional forming required to secure the formed tubular bar in an intended location. The bar may be tubular, having a substantially circular outer surface and inner surface, although other shapes and cross-sections are contemplated as well. As used herein, the term “tubular” may refer to any elongate, hollow shape, which need not be circular or cylindrical.
The terms “approximately”, “about”, and “substantially” as used herein represent an amount equal to or close to a stated amount that still performs a desired function or achieves a desired result. For example, the terms “approximately,” “about,” and “substantially” may refer to an amount that is within less than 10% of, within less than 5% of, within less than 1% of, within less than 0.1% of, and within less than 0.01% of the stated amount.
The term “room temperature” as used herein has its ordinary meaning as known to those skilled in the art and may include temperatures within the range of about 16° C. (60° F.) to about 32° C. (90° F.).
In general, embodiments of the present disclosure comprise low-alloy carbon steel pipes and methods of manufacture. As discussed in greater detail below, through a combination of steel composition and heat treatment, a steel pipes may be provided having a final microstructure that provides selected mechanical properties of interest, including one or more of minimum yield strength, toughness, hardness, and corrosion resistance in high wall thickness pipes (e.g., WT greater than or equal to about 35 mm).
Embodiments of the disclosed steel compositions may comprise not only carbon (C) but also manganese (Mn), silicon (Si), chromium (Cr), nickel (Ni), molybdenum (Mo), vanadium (V), aluminum (Al), nitrogen (N), and calcium (Ca). Additionally, one or more of the following elements may be optionally present and/or added as well: tungsten (W), niobium (Nb), titanium (Ti), boron (B), zirconium (Zr), and tantalum (Ta). The remainder of the composition may comprise iron (Fe) and impurities. In certain embodiments, the concentration of impurities may be reduced to as low an amount as possible. Embodiments of impurities may include, but are not limited to, copper (Cu), sulfur (S), phosphorous (P), arsenic (As), antimony (Sb), tin (Sn), bismuth (Bi), oxygen (O), and hydrogen (H).
For example, embodiments of the low-alloy steel composition may comprise (in weight % unless otherwise noted):
Carbon within the range between about 0.05% to about 0.16%;
Manganese within the range between about 0.20% to about 0.90%;
Silicon within the range between about 0.10% to about 0.50%;
Chromium within the range between about 1.20% to about 2.60%;
Nickel within the range between about 0.050% to about 0.50%;
Molybdenum within the range between about 0.80% to about 1.20%;
Tungsten less than or equal to about 0.80%
Niobium less than or equal to about 0.030%;
Titanium less than or equal to about 0.020%;
Vanadium within the range between about 0.005% to about 0.12%;
Aluminum within the range between about 0.008% to about 0.040%;
Nitrogen within the range between about 0.0030% to about 0.012%;
Copper less than or equal to about 0.3%;
Sulfur less than or equal to about 0.01%;
Phosphorous less than or equal to about 0.02%;
Calcium within the range between about 0.001%-0.005%;
Boron less than or equal to about 0.0020%;
Arsenic less than or equal to about 0.020%;
Antimony less than or equal to about 0.005%;
Tin less than or equal to about 0.020%;
Zirconium less than or equal to 0.03%;
Tantalum less than or equal to 0.03%;
Bismuth less than about 0.0050%;
Oxygen less than about 0.0030%;
Hydrogen less than or equal to about 0.00030; and
the balance of the composition comprising iron and impurities.
Heat treatment operations performed on pipes formed from the steel composition may include, but are not limited to, quenching and tempering (Q+T). The quenching operation may include reheating a pipe from about room temperature, after hot forming, to a temperature that austenitizes the pipe, followed by a rapid quench. For example, a pipe may be heated to a temperature within the range between about 900° C. to about 1060° C. and held at about the austenitizing temperature for a selected soaking time. Cooling rates during the quench may be selected so as to achieve a selected cooling rate at about the mid-wall of the pipe. For example, pipes may be cooled so as to achieve cooling rates greater than or equal to about 7° C./s at the mid-wall.
Quenching pipes having a WT greater than or equal to about 35 mm and the composition described above may promote the formation of a volume percent of martensite greater than about 50%, preferably greater than about 70% and more preferably greater than about 90% within the pipe. The remaining microstructure of the pipe may comprise lower bainite, with substantially no ferrite, upper bainite, or granular bainite.
Following the quenching operations, pipes may be further subjected to tempering. Tempering may be conducted at a temperature within the range between about 680° C. to about 760° C., depending upon the composition of the steel and the target yield strength. In addition to martensite and lower bainite, the microstructure may further exhibit an average prior austenite grain size, measured according to ASTM E112, of about 15 μm or about 20 μm to about 100 μm. The microstructure may also exhibit an average packet size of less than about 6 μm. The microstructure may further exhibit one or more precipitates. For example, precipitates of a first type given by any of MX, M2X, where M=V, Mo, Nb, or Cr and X═C or N may be present in the microstructure. Precipitates of the first type may have an average diameter less than or equal to about 40 nm. In certain embodiments, precipitates of a second type given by any of M3C, M6C, and M23C6 may be present. Precipitates of the second type may have an average diameter between about 80 nm to about 400 nm.
In one embodiment, a steel pipe having a WT greater than about 35 mm and the composition and microstructure discussed above may possess the following properties:
In another embodiment, a steel pipe having a WT greater than about 35 mm may be formed having the following properties:
In each of the above embodiments, formed pipe may further exhibit the following impact and hardness properties:
In each of the above embodiments, formed pipe may further exhibit the following resistance to sulfide stress corrosion (SSC) cracking and hydrogen induced cracking (MC). SSC testing is conducted according to NACE TM 0177 using solution A with a test duration of about 720 hours. HIC testing is conducted according to NACE TM 0284-2003 Item 21215 using NACE solution A and test duration of about 96 hours:
HIC
SSC
With reference to
Operation 102 of the method 100 preferably comprises fabrication of the steel and production of a solid metal billet capable of being pierced and rolled to form a metallic tubular bar. In further embodiments, selected steel scrap, cast iron, and sponge iron may be employed to prepare the raw material for the steel composition. It may be understood, however, that other sources of iron and/or steel may be employed for preparation of the steel composition.
Primary steelmaking may be performed using an electric arc furnace to melt the steel, decrease phosphorous and other impurities, and achieve a selected temperature. Tapping and deoxidation, as well as addition of alloying elements may be further performed.
One of the main objectives of the steelmaking process is to refine the iron by removal of impurities. In particular, sulfur and phosphorous are prejudicial for steel because they degrade the mechanical properties of the steel. In one embodiment, secondary steelmaking may be performed in a ladle furnace and trimming station after primary steelmaking to perform specific purification steps.
During these operations, very low sulfur contents may be achieved within the steel, calcium inclusion treatment is performed, and inclusion flotation is performed. In one embodiment inclusion flotation may be performed by bubbling inert gases in the ladle furnace to force inclusions and impurities to float. This technique produces a fluid slag capable of absorbing impurities and inclusions. In this manner, a high quality steel having the desired composition with a low inclusion content may be provided.
Table 1 illustrates embodiments of the steel composition, in weight percent (wt. %) unless otherwise noted.
Carbon (C) is an element whose addition to the steel composition may inexpensively raise the strength of the steel and refine the microstructure, reducing the transformation temperatures. In an embodiment, if the C content of the steel composition is less than about 0.05%, it may be difficult in some embodiments to obtain the strength desired in articles of manufacture, particularly tubular products. On the other hand, in other embodiments, if the steel composition has a C content greater than about 0.16%, in some embodiments, toughness is impaired, and weldability may decrease, making more difficult and expensive any welding process if joining is not performed by thread joints. In addition, the risk of developing quenching cracks in steels with high hardenability increases with the carbon content. Therefore, in an embodiment, the C content of the steel composition may be selected within the range between about 0.05% to about 0.16%, preferably within the range between about 0.07% to about 0.14%, and more preferably within the range between about 0.08% to about 0.12%.
Manganese (Mn) is an element whose addition to the steel composition may be effective in increasing the hardenability, strength and toughness of the steel. In an embodiment, if the Mn content of the steel composition is less than about 0.20% it may be difficult in some embodiments to obtain the desired strength in the steel. However, in another embodiment, if the Mn content of the steel composition exceeds about 0.90%, in some embodiments banding structures may become marked in some embodiments, and toughness and HIC/SSC resistance may decrease. Therefore, in an embodiment, the Mn content of the steel composition may be selected within the range between about 0.20% to about 0.90%, preferably within the range between about 0.30% to about 0.60%, and more preferably within the range between about 0.30% to about 0.50%.
Silicon (Si) is an element whose addition to the steel composition may have a deoxidizing effect during steel making process and may also raise the strength of the steel (e.g., solid solution strengthening). In an embodiment, if the Si content of the steel composition is less than about 0.10%, the steel in some embodiments may be poorly deoxidized during steelmaking process and exhibit a high level of micro-inclusions. In another embodiment, if the Si content of the steel composition exceeds about 0.50%, both toughness and formability of the steel may decrease in some embodiments. Si content higher than about 0.5% are also recognized to have a detrimental effect on surface quality when the steel is processed at high temperatures (e.g., temperatures greater than about 1000° C.) in oxidizing atmospheres, because surface oxide (scale) adherence is increased due to fayalite formation and the risk of surface defect is higher. Therefore, in an embodiment, the Si content of the steel composition may be selected within the range between about 0.10% to about 0.50%, preferably within the range between about 0.10% to about 0.40%, and more preferably within the range between about 0.10% to about 0.25%.
Chromium (Cr) is an element whose addition to the steel composition may increase hardenability, decrease transformation temperatures, and increase tempering resistance of the steel. Therefore the addition of Cr to steel compositions may be desirable for achieving high strength and toughness levels. In an embodiment, if the Cr content of the steel composition is less than about 1.2%, it may be difficult in to obtain the desired strength and toughness, some embodiments. In another embodiment, if the Cr content of the steel composition exceeds about 2.6%, the cost may be excessive and toughness may decrease due to enhanced precipitation of coarse carbides at grain boundaries, in some embodiments. In addition, weldability of the resultant steel may be reduced, making the welding process more difficult and expensive, if joining is not performed by thread joints. Therefore, in an embodiment, the Cr content of the steel composition may be selected within the range between about 1.2% to about 2.6%, preferably within the range between about 1.8% to about 2.5%, and more preferably within the range between about 2.1% to about 2.4%.
Nickel (Ni) is an element whose addition may increase the strength and toughness of the steel. However, in an embodiment, when Ni addition exceeds about 0.5%, a negative effect on scale adherence has been observed, with higher risk of surface defect formation. Also, in other embodiments, Ni contents higher than about 1% are recognized to have a detrimental effect on sulfide stress corrosion cracking. Therefore, in an embodiment, the Ni content of the steel composition may be selected within the range between about 0.05% to about 0.5%.
Molybdenum (Mo) is an element whose addition to the steel composition may improve hardenability and hardening by solid solution and fine precipitation. Mo may assist in retarding softening during tempering, promoting the formation of very fine MC and M2C precipitates. These particles may be substantially uniformly distributed in the matrix and may also act as beneficial hydrogen traps, slowing down the atomic hydrogen diffusion towards the dangerous traps, usually at grain boundaries, which behave as crack nucleation sites. Mo also reduces the segregation of phosphorous to grain boundaries, improving resistance to inter-granular fracture, with beneficial effects also on SSC resistance because high strength steels which suffer hydrogen embrittlement exhibit an intergranular fracture morphology. Therefore, by increasing the Mo content of the steel composition, the desired strength can be achieved at higher tempering temperatures, which promote better toughness levels. In an embodiment, in order to exert the effect thereof, the Mo content may be greater than or equal to about 0.80%. However, in other embodiments, for Mo contents higher than about 1.2% a saturation effect on hardenability is noted and weldability may be reduced. As Mo ferroalloy is expensive, in an embodiment, the Mo content of the steel composition may be selected within the range between about 0.8 to about 1.2%, preferably within the range between about 0.9% to about 1.1%, and more preferably within the range between about 0.95% to about 1.1%.
Tungsten (W) is an element whose addition to the steel composition is optional and may increase the strength at room and elevated temperatures by forming tungsten carbide which develops secondary hardening. W is preferably added when the steel use is required at high temperatures. The behavior of W is similar to that of Mo in terms of hardenability but its effectiveness is about one half of that of Mo. Tungsten reduces the steel oxidation and, as a result, less scale is formed during reheating processes at high temperatures. However, as its cost is very high, in an embodiment, the W content of the steel composition may selected to be less than or equal to about 0.8%.
Niobium (Nb) is an element whose addition to the steel composition is optional and may be provided to form carbides and nitrides and may be further used to refine the austenitic grain size during hot rolling and re-heating before quenching. However Nb is not needed in embodiments of present steel composition to refine the austenite grains as a predominant martensite structure is formed and a fine packet is formed even in the case of coarse austenite grains when low transformation temperatures are promoted through a proper balance of other chemical elements such as Cr, Mo, and C.
Nb precipitates as carbonitride may increase the steel strength by particle dispersion hardening. These fine and round particles may be substantially uniformly distributed in the matrix and also act as hydrogen traps, beneficially slowing down the atomic hydrogen diffusion towards the dangerous traps, usually at grain boundaries, which behave as crack nucleation sites. In an embodiment, if the Nb content is higher than about 0.030%, a coarse precipitate distribution that impair toughness may be formed. Therefore, in an embodiment, the Nb content of the steel composition may be selected to be less than or equal to about 0.030%, preferably less than or equal to about 0.015%, and more preferably less than or equal to about 0.01%.
Titanium (Ti) is an element whose addition to the steel composition is optional and may be provided to refine austenitic grain size in high temperature processes, forming nitrides and carbonitrides. However it is not needed in embodiments of present steel composition, except when it is used to protect boron that remains in solid solution improving hardenability, especially in the case of pipes with wall thickness greater than about 25 mm. For example, Ti binds nitrogen and avoids BN formation). Additionally, in certain embodiments, when Ti is present in concentrations higher than about 0.02%, coarse TiN particles may be formed that impair toughness. Accordingly, in an embodiment, the Ti content of the steel composition may be less than or equal to about 0.02%, and more preferably less than or equal to about 0.01% when boron is below about 0.0010%.
Vanadium (V) is an element whose addition to the steel composition may increase strength by carbonitride precipitation during tempering. These fine and round particles may also be substantially uniformly distributed within the matrix and act as beneficial hydrogen traps. In an embodiment, if the V content is less than about 0.05%, it may be in some embodiments difficult to obtain the desired strength. However, in another embodiment, if the V content is higher than 0.12%, a large volume fraction of vanadium carbide particles may be formed with subsequent reduction in toughness. Therefore, in certain embodiments, the Nb content of the steel composition may be selected to be less than or equal to about 0.12%, preferably within the range between about 0.05% to about 0.10%, and more preferably within the range between about 0.05% to about 0.07%.
Aluminum (Al) is an element whose addition to the steel composition has a deoxidizing effect during steel making process and may refine the steel grain. In an embodiment, if the Al content of the steel composition is higher than about 0.040%, coarse precipitates of AlN that impair toughness and/or Al-rich oxides (e.g., non-metallic inclusions) that impair HIC and SSC resistance may be formed. Accordingly, in an embodiment, the Al content of the steel may be selected to be less than or equal to about 0.04%, preferably less than or equal to about 0.03%, and more preferably less than or equal to about 0.025%.
Nitrogen (N) is an element whose content within the steel composition is preferably selected to be greater than or equal to about 0.0030%, in one embodiment, in order to form carbonitrides of V, Nb, Mo and Ti. However, in other embodiments, if the N content of the steel composition exceeds about 0.0120%, the toughness of the steel may be degraded. Therefore, the N content of the steel composition may be selected within the range between about 0.0030% to about 0.0120%, preferably within the range between about 0.0030% to about 0.0100%, and more preferably within the range between about 0.0030% to about 0.0080%.
Copper (Cu) is an impurity element that is not needed in embodiments of the steel composition. However, depending on the manufacturing process, the presence of Cu may be unavoidable. Therefore, the Cu content may be limited to as low as possible. For example, in an embodiment, the Cu content of the steel composition may be less than or equal to about 0.3%, preferably less than or equal to about 0.20%, and more preferably less than or equal to about 0.15%.
Sulfur (S) is an impurity element that may decrease both toughness and workability of the steel, as well as HIC/SSC resistance. Accordingly, the S content of the steel in some embodiments may be kept as low as possible. For example, in an embodiment, the Cu content of the steel composition may be less than or equal to about 0.01%, preferably less than or equal to about 0.005%, and more preferably less than or equal to about 0.003%.
Phosphorous (P) is an impurity element that may cause the toughness and HIC/SSC resistance of high strength steel to decrease. Accordingly, the P content in some embodiments may be kept as low as possible. For example, in an embodiment, the P content of the steel composition may be less than or equal to about 0.02%, preferably less than or equal to about 0.012%, and more preferably less than or equal to about 0.010%.
Calcium (Ca) is an element whose addition to the steel composition may assist with control of the shape of inclusions and enhancement of the HIC resistance by forming fine and substantially round sulfides. In an embodiment, in order to provide these benefits, the Ca content of the steel composition may be selected to be greater than or equal to about 0.0010% when the sulfur content of the steel composition is higher than about 0.0020%. However in other embodiments, if the Ca content of the steel composition exceeds about 0.0050% the effect of the Ca addition may be saturated and the risk of forming clusters of Ca-rich non-metallic inclusions that reduce HIC and SSC resistance may be increased. Accordingly, in certain embodiments, the maximum Ca content of the steel composition may be selected to be less than or equal to about 0.0050%, and more preferably less than or equal to about 0.0030%, while the minimum Ca content may be selected to be greater than or equal to about 0.0010%, and most preferably to greater than or equal to about 0.0015%.
Boron (B) is an element whose addition to the steel composition is optional and may be provided for improving the hardenability of the steel. B can be used for inhibiting ferrite formation. In an embodiment, the lower limit of the B content of the steel composition to provide these beneficial effects may be about 0.0005%, while the beneficial effects may be saturated with boron contents higher than about 0.0020%. Therefore, in selected embodiments, the maximum B content of the steel composition may be selected to be less than or equal to about 0.0020%.
Arsenic (As), tin (Sn), antimony (Sb) and bismuth (Bi) are impurity elements that are not needed in embodiments of the steel composition. However, depending on the manufacturing process, the presence of these impurity elements may be unavoidable. Therefore, the As and Sn contents within the steel composition may be selected to be less than or equal to about 0.020%, and more preferably less than or equal to about 0.015%. The Sb and Bi contents may be selected to be less than or equal to about 0.0050%.
Zirconium (Zr) and tantalum (Ta) are elements that act as strong carbide and nitride formers, similar to Nb and Ti. These elements may be optionally added to the steel composition, as they are not needed in embodiments of present steel composition to refine the austenite grains. Zr and Ta fine carbonitrides may increase the steel strength by particle dispersion hardening and may also act as beneficial hydrogen traps, slowing down the atomic hydrogen diffusion towards the dangerous traps. In an embodiment, if the Zr or Ta content is greater than or equal to about 0.030%, a coarse precipitate distribution that may impair toughness of the steel may be formed. Zirconium may also act as a deoxidizing element in steel and combine with the sulfur. However, as addition to steel in order to promote globular non-metallic inclusions, in certain embodiments, Ca may be preferred over Zr. Therefore, the content of Zr and Ta within the steel composition may be selected to be less than or equal to about 0.03%.
The total oxygen (O) content of the steel composition is the sum of the soluble oxygen and the oxygen in the non-metallic inclusions (oxides). As it is practically the oxygen content in the oxides in a well deoxidized steel, an oxygen content that is too high means a high volume fraction of non-metallic inclusions and less resistance to HIC and SSC. Accordingly, in an embodiment, the oxygen content of the steel may be selected to be less than or equal to about 0.0030%, preferably less than or equal to about 0.0020%, and more preferably less than or equal to about 0.0015%.
Following the production of the fluid slag having a composition as described above, the steel may be cast into a round solid billet having a substantially uniform diameter along the steel axis. For example, round billets having a diameter within the range between about 330 mm to about 420 mm may be produced in this manner.
The billet thus fabricated may be formed into a tubular bar through hot forming processes 104. In an embodiment, a solid, cylindrical billet of clean steel may be heated to a temperature of about 1200° C. to 1.340° C., preferably about 1280° C. For example, the billet may be reheated by a rotary heath furnace. The billet may be further subject to a rolling mill. Within the rolling mill, the billet may be pierced, in certain preferred embodiments utilizing the Manessmann process, and hot rolling is used to substantially reduce the outside diameter and wall thickness of the tube, while the length is substantially increased. In certain embodiments, the Manessmann process may be performed at temperatures within the range between about 1200° C. to about 1280° C. The obtained hollow bars may be further hot rolled at temperatures within the range between about 1000° C. to about 1200° C. in a retained mandrel continuous mill. Accurate sizing may be carried out by a sizing mill and the seamless tubes cooled in air to about room temperature in a cooling bed. For example, pipes with outer diameters (OD) within the range between about 6 inches to about 16 inches may be formed in this manner.
After rolling the pipes may be in-line heated, without cooling at room temperature, by an intermediate furnace for making temperature more uniform, and accurate sizing may be carried out by a sizing mill. Subsequently, the seamless pipes may be cooled in air down to room temperature in a cooling bed. In the case of a pipe having a final OD greater than about 16 inches, the pipes produced by the medium size mill may be processed by a rotary expansion mill. For example, medium size pipes may be reheated by a walking beam furnace to a temperature within the range between about 1150° C. to about 1250° C., expanded to the desired diameter by the expander mill at a temperature within the range between about 1100° C. to about 1200° C., and in-line reheated before final sizing.
In a non-limiting example, a solid bar may be hot formed as discussed above into a tube possessing an outer diameter within the range between about 6 inches to about 16 inches and a wall thickness greater than about 35 mm.
The final microstructure of the formed pipe may be determined by the composition of the steel provided in operation 102 and heat treatments performed in operations 106. The composition and microstructure, in turn, may give rise to the properties of the formed pipe.
In one embodiment, promotion of martensite formation may refine the packet size (the size of the regions separated by high-angle boundaries that offer higher resistance to crack propagation; the higher the misorientation, the higher the energy a crack requires to cross the boundary) and improve the toughness of the steel pipe for a given yield strength. Increasing the amount of martensite in as-quenched pipes may further allow the use of higher tempering temperatures for a given strength level. Therefore, in an embodiment, it is a goal of the method to achieve a predominantly martensitic microstructure at relatively low temperatures (e.g., transformation of austenite at temperatures less than or equal to about 450° C. In an embodiment, the martensitic microstructure may comprise a volume percent of martensite greater than or equal to about 50%. In further embodiments, the volume percent of martensite may be greater than or equal to about 70%. In further embodiments, the volume percent of martensite may be greater than or equal to about 90%.
In another embodiment, hardenability of the steel, the relative ability of the steel to form martensite when quenched, may be improved through the composition and microstructure. In one aspect, addition of elements such as Cr and Mo are effective in reducing the transformation temperature of martensite and bainite and increase the resistance to tempering. Beneficially, a higher tempering temperature may then be used to achieve a given strength level (e.g., yield strength). In another aspect, a relatively coarse prior austenite grain size (e.g., about 15 or 20 μm to about 100 μm) may improve hardenability.
In a further embodiment, the sulfide stress corrosion cracking (SSC) resistance of the steel may be improved through the composition and microstructure. In one aspect, the SSC may be improved by increased content of martensite within the pipe. In another aspect, tempering at very high temperatures may improve the SSC of the pipe, as discussed in greater detail below.
In order to promote martensite formation at temperatures less than or equal to about 450° C., the steel composition may further satisfy Equation 1, where the amounts of each element are given in wt. %:
60C %+Mo %+1.7Cr %>10 Eq. 1
If a significant amount of bainite (e.g., less than about 50 volume %) is present after quenching, the temperature at which the bainite forms should be less than or equal to about 540° C. in order to promote a relatively fine packet, with substantially no upper bainite or granular bainite (a mixture of bainitic dislocated-ferrite and islands of high C martensite and retained austenite).
In order to promote the bainite formation at a temperature less than or equal to about 540° C. (e.g., lower bainite), the steel composition may additionally satisfy Equation 2, where the amounts of each element are given in wt. %:
60C %+41Mo %+34Cr %>70 Eq. 2
Notably, normalizing (e.g., austenitizing followed by cooling in still air), may not achieve the desired martensite microstructure because the typical average cooling rates between about 800° C. and 500° C. for pipes of wall thickness between about 35 mm and about 60 mm is lower than about 1° C./s. Water quenching may be employed to achieve the desired cooling rates at about the pipe mid-wall and form martensite and lower bainite at temperatures lower than about 450° C. and about 540° C., respectively. Therefore, the as-rolled pipes may be reheated in a furnace and water quenched in quenching operation 106A after air-cooling from hot rolling.
For example, in one embodiment of the austenitizing operations 106A, the temperatures of the zones of the furnace may be selected in order to allow the pipe to achieve the target austenitizing temperature with a tolerance lower than about +/−20° C. Target austenitizing temperatures may be selected within the range between about 900° C. to about 1060° C. The heating rate may be selected within the range between about 0.1° C./s to about 0.2° C./s. The soaking time, the time from when the pipe achieves the final target temperature minus about 10° C. and the exit from the furnace, may be selected within the range between about 300 sec to about 1800 sec. Austenitizing temperatures and holding times may be selected depending on chemical composition, wall thickness, and desired austenite grain size. At the exit of the furnace, the pipe may be descaled to remove the surface oxide and is rapidly moved to a water quenching system.
In the quenching operations 106B, external and internal cooling may be employed to achieve the desired cooling rates at about the mid-wall of the pipe (e.g., greater than about 7° C./s). As discussed above, cooling rates within this range may promote the formation of a volume percent of martensite greater than about 50%, preferably greater than about 70%, and more preferably greater than about 90%. The remaining microstructure may comprise lower bainite, (i.e. bainite formed at temperatures lower than about 540° C. with a typical morphology including fine precipitation within the bainite laths, without coarse precipitates at lath boundaries as in the case of upper bainite, which is usually formed at temperatures higher than about 540° C.).
In one embodiment, the water quench of quenching operations 106B may be performed by dipping the pipe in a tank containing stirred water. The pipe may be rapidly rotated during quenching to make the heat transfer high and uniform and avoid pipe distortion. Additionally, in order to remove the steam developed inside the pipe, an inner water jet may also be employed. In certain embodiments, the water temperature may not be higher than about 40° C., preferably less than about 30° C. during quenching operations 106B.
After quenching operations 106B, the pipe may be introduced in another furnace for the tempering operations 106C. In certain embodiments, the tempering temperature may be selected to be sufficiently high so as to produce a relatively low dislocation density matrix and more carbides with a substantially round shape (i.e., a higher degree of spheroidization). This spheroidization improves the impact toughness of the pipes, as needle shaped carbides at lath and grain boundaries may provide easier crack paths.
Tempering the martensite at temperatures sufficiently high to produce more spherical, dispersed carbides may promote trans-granular cracking and better SSC resistance. Crack propagation may be slower in steels that possess a high number of hydrogen trapping sites and fine, dispersed precipitates having spherical morphologies give better results.
By forming a microstructure including tempered martensite, as opposed to a banded microstructure (e.g., ferrite-pearlite or ferrite-bainite), the HIC resistance of the steel pipe may be further increased.
In one embodiment, the tempering temperature may be selected within the range between about 680° C. to about 760° C. depending on the chemical composition of the steel and the target yield strength. The tolerances for the selected tempering temperature may be within the range of about ±15° C. The pipe may be heated at a rate between about 0.1° C./s to about 0.2° C./s to the selected tempering temperature. The pipe may be further held at the selected tempering temperature for a duration of time selected within the range between about 1800 sec to about 5400 sec.
Notably, the packet size is not significantly influenced by the tempering operations 106C. However, packet size may decrease with a reduction of the temperature at which austenite transforms. In traditional low-carbon steels with carbon equivalents lower than about 0.43%, tempered bainite may show a coarser packet size (e.g., about 7 μm to about 12 μm) as compared with that of the tempered martensite within the instant application (e.g. less than or equal to about 6 μm, such as from within the range about 6 μm to about 2 μm).
The martensite packet size is nearly independent of the average austenite grain size and may remain fine (e.g., an average size less than or equal to about 6 μm) even in the case of relatively coarse average austenite grain size (e.g., about 15 μm or about 20 μm to about 100 μm).
Finishing operations 110 may include, but are not limited to, straightening and bending operations. Straightening may be performed at temperatures below the tempering temperature and above about 450° C.
In one embodiment, bending may be performed by hot induction bending. Hot induction bending is a hot deformation process which concentrates in a narrow zone, referred to as hot tape, that is defined by an induction coil (e.g., a heating ring) and a quenching ring that sprays water on the external surface of the structure to be bent. A straight (mother) pipe is pushed from its back, while the front of the pipe is clamped to an arm constrained to describe a circular path. This constraint provokes a bending moment on the entire structure, but the pipe is plastically deformed substantially only within correspondence of the hot tape. The quenching ring plays therefore two simultaneous roles: (i) to define the zone under plastic deformation and (ii) to in-line quench the hot bend.
The diameter of both heating and quenching rings is about 20 mm to about 60 mm larger than the outside diameter (OD) of the mother pipe. The bending temperature at both exterior and interior surfaces of the pipe may be continuously measured by pyrometers.
In conventional pipe fabrication, the bends may be subjected to a stress relieving treatment after bending and quenching by a tempering treatment at a relatively low temperature to achieve the final mechanical properties. However, it is recognized that the in-line quenching and tempering operations performed during finishing operations 110 may produce a microstructure that is different than that obtained from the off-line quenching and tempering operations 106B, 106C. Therefore, in an embodiment of the disclosure, as discussed above in operations 106B, 106C, in order to substantially regenerate the microstructure obtained after operations 106B, 106C. Therefore, the bends may, be reheated in a furnace and then rapidly immersed into a quenching tank with stirred water and then tempered in a furnace.
In an embodiment, the temper after bending may be performed at a temperature within the range between about 710° C. to about 760° C. The pipe may be heated at a rate within the range between about 0.05 C/s to about 0.2° C./s. A hold time within the range between about 1800 sec to about 5400 sec may be employed after the target tempering temperature has been achieved.
Even when the austenite grain is coarse, as it is in this instance, the packet size of the steel after quenching and tempering may be maintained below approximately 6 μm if a predominant martensite structure (e.g., martensite greater than about 50% in volume) is formed and lower bainite forms at relatively low temperatures (<540° C.).
Packet size is measured as average lineal intercept on images taken by Scanning Electron Microscopy (SEM) using the Electron Back Scattered Diffraction (EBSD) signal, and considering high-angle boundaries those with misorientation greater than about 45°. Measurement by the lineal intercept method gave distribution shown in
On the quenched and tempered pipe, fine precipitates of a first type given by any of MX, M2X, where M is Mo or Cr or V, Nb, Ti when present, and X is C or N, with size (e.g., average diameter) less than about 40 nm were also detected by Transmission Electron Microscopy (TEM), in addition to coarse precipitates of the type M3C, M6C, M23C6 with an average diameter within the range between about 80 nm to about 400 nm.
The total volume percentage of non-metallic inclusions was below about 0.05%, preferably below about 0.04%. The number of inclusions per square mm of examined area of oxides with size larger than about 15 μm was below about 0.4/mm2. Substantially only modified round sulfides were present.
In the following examples, the microstructural and mechanical properties and impact of steel pipes formed using embodiments of the steel making method discussed above are discussed. In particular, microstructural parameters including austenite grain size, packet size, martensite volume, lower bainite volume, volume of non-metallic inclusions, and inclusions of greater than about 15 μm were examined for embodiments of the compositions and heat treatment conditions discussed above. Corresponding mechanical properties, including yield and tensile strengths, hardness, elongation, toughness, and HIC/SSC resistance are further discussed.
The microstructural and mechanical properties of the steel of Table 2 were investigated. With respect to the measurement of microstructural parameters, austenite grain size (AGS) was measured in accordance with ASTM E112, packet size was measured using an average lineal intercept on images taken by scanning electron microscopy (SEM) using the electron backscatter diffraction (EBSD) signal, the volume of martensite was measured in accordance with ASTM E562, the volume of lower bainite was measured in accordance with ASTM E562, the volume percentage of non-metallic inclusions was measured by automatic image analysis using optical microscopy in accordance with ASTM E1245, and the presence of precipitates was investigated by transmission electron microscopy (TEM) using the extraction replica method.
With respect to the mechanical properties, yield strength, tensile strength, and elongation were measured in accordance with ASTM E8, hardness was measured in accordance with ASTM E92, impact energy was evaluated on transverse Charpy V-notch specimens according to ISO 148-1, ductile-to-brittle transition temperature was evaluated on transverse Charpy V-notch specimens in accordance with ASTM E208, crack tip opening displacement was measured according to BS7488 part 1 at about −60° C., HIC evaluation was performed in accordance with NACE Standard TM0284-2003, Item No. 21215 using NACE solution A and a test duration of about 96 hours. SSC evaluation was performed in accordance with NACE TM0177 using test solution A and conducted for a test duration of about 720 hours at about 90% yield stress.
A heat of about 90 t, with the chemical composition range shown in Table 2, was manufactured by electric arc furnace.
After tapping, deoxidation, and alloying additions, secondary metallurgy operations were carried out in, a ladle furnace and trimming station. After calcium treatment and vacuum degassing, the liquid steel was then continuously cast on a vertical casting machine as round bars of approximately 330 mm diameter.
The as-cast bars were re-heated by a rotary heath furnace to a temperature of about 1300° C., hot pierced, and the hollows were hot rolled by a retained mandrel multi-stand pipe mill and subjected to hot sizing in accordance process described above with respect to
The as-rolled pipes were subsequently austenitized by heating to a temperature of about 920° C. for approximately 5400 sec by a walking beam furnace, descaled by high pressure water nozzles, and externally and internally water quenched using a tank with stirred water and an inner water nozzle. The austenitizing heating rate was approximately 0.16° C./s. The cooling rate employed during quenching was approximately 15° C./s. The quenched pipes were rapidly moved to another walking beam furnace for tempering treatment at a temperature of about 740° C. for a total time of about 9000 sec and a soaking time of about 4200 sec. The tempering heating rate was approximately 0.12° C./s. The cooling rate employed during tempering was approximately less than about 0.1° C./s. All the quenched and tempered (Q&T) pipes were hot straightened.
The main parameters characterizing the microstructure and non-metallic inclusions of the pipes of Example 1 are shown in Table 4.
The mechanical properties of the pipes of Example 1 are shown in Tables 5, 6, and 7. Table 5 presents the tensile, elongation, hardness, and toughness properties of the quenched and tempered pipes. Table 6 presents the yield strength, fracture appearance transition temperature, crack tip opening displacement, and ductility transition temperature after a simulated post-weld heat treatment. The post-weld heat treatment comprised heating and cooling at a rate of about 80° C./h to a temperature of about 690° C. with a soaking times of about 5 h. Table 7 presents the measured HIC and SSC resistance of the quenched and tempered pipes.
It was found from the testing results above (Table 5, Table 6, and Table 7) that the quenched and tempered pipes are suitable to develop a 65 ksi grade, characterized by:
The quenched and tempered pipes of Example 1 were used to manufacture bends having a radius of approximately 5 times the outer diameter of the pipe (5 D).
The pipes were subjected to hot induction bending by heating to a temperature of approximately 850° C.+/−25° C. and in-line water quenching. The bends were then reheated to a temperature of about 920° C. for approximately 15 min holding in a car furnace, moved to a water tank, and immersed in stirred water. The minimum temperature of the bends was higher than about 860° C. just before immersion in the water tank and the temperature of the water of the tank was maintained below approximately 40° C. The microstructure of the as-quenched bend at about the mid-wall of the pipe is illustrated in
Following the quenching operation, the as-quenched bends were tempered in a furnace set at a temperature of about 730° C. using an approximately 40 min holding time.
It was found from the testing results above (Table 8, Table 9) that the quenched and tempered pipes are suitable to develop a 70 ksi grade, characterized by:
In this comparative example, quenched and tempered pipes having an outer diameter of about 219.1 mm and wall thickness of about 44 mm, made of a typical line pipe steel with a low carbon equivalent of 0.4% (Table 10), were used to manufacture hot induction bends, off-line quench and temper, using embodiments of the process previously described.
The produced seamless pipes, were austenitized at about 920° C. using a soaking time of about 600 sec, as discussed above, by a walking beam furnace. The pipes were further descaled by high pressure water nozzles and externally and internally water quenched using a tank with stirred water and an inner water nozzle. The quenched pipes were rapidly moved to another walking beam furnace for tempering treatment at about 660° C. to about 670° C. Each of the quenched and tempered pipes were hot straightened.
The Q&T pipes were further subjected to hot induction bending by heating to a temperature of about 850° C.+/−25° C. and in-line water quenched. The bends were then reheated at about 920° C. for an approximately 30 min hold time in a car furnace, moved to a water tank and immersed in stirred water. The minimum temperature of the bends was greater than about 860° C. just before immersion in the water tank and the temperature of the water of the tank was maintained below about 40° C. The microstructure at about the mid-wall of the as-quenched bend is illustrated in
A predominant microstructure within the as-quenched pipe was granular bainite (a mixture of bainitic dislocated-ferrite and islands of high C martensite and retained austenite, MA constituent), which is significantly different from that of the high Cr-high Mo steel in
The as-quenched bends were further tempered in a furnace set at about 670° C. using an approximately 30 min holding time.
The main parameters which characterize the microstructure and non-metallic inclusions of the Q&T bends are shown in Table 11.
From the forgoing, it may be observed that pipes having quenched and tempered bends, as they are manufactured with a steel that does not develop enough hardenability, exhibit a predominant granular bainite microstructure. Moreover, the packet size is larger than that of Example 2.
Moreover, while these quenched and tempered bends are able to achieve the minimum yield strength of 450 MPa, i.e. grade X65 (Table 12), they have a lower toughness with higher transition temperatures and a lower resistance to SSC, as compared to Example 2, due to their different microstructure.
Although the foregoing description has shown, described, and pointed out the fundamental novel features of the present teachings, it will be understood that various omissions, substitutions, and changes in the form of the detail of the apparatus as illustrated, as well as the uses thereof, may be made by those skilled in the art, without departing from the scope of the present teachings. Consequently, the scope of the present teachings should not be limited to the foregoing discussion, but should be defined by the appended claims.
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| Number | Date | Country | |
|---|---|---|---|
| 20120204994 A1 | Aug 2012 | US |
