High strength steel having good toughness

This application is related to Applicant's co-pending application entitled ULTRA HIGH STRENGTH STEEL HAVING GOOD TOUGHNESS, Ser. No. 13/031,133, filed Feb. 18, 2011, the entirety of which is hereby incorporated by reference.

1. Field of the Invention

The present invention relates generally to metal production and, in certain embodiments, relates to methods of producing metallic tubular bars having high strength while concurrently possessing good toughness.

2. Description of the Related Art

Seamless steel tubes are widely used in a variety of industrial applications. Due to requirements for higher load bearing capacity, situations of dynamic stresses, and the need for lighter components, there is an increasing demand for the development of steel tubes possessing increased strength and toughness.

In the oil industry, perforating guns comprising steel tubes containing explosive charges are used to deliver explosive charges to selected locations of wells. The steel tubes used as perforating gun carriers are subjected to very high external collapse loads that are exerted by the hydrostatic well pressure. On the other hand, during detonation, the steel tubes are also subjected to very high dynamic loads. To address this issue, efforts have been directed to the development of steel tubes with high strength, while at the same time maintaining very good impact toughness.

At present, the highest available steel grade in the market has a minimum yield strength of about 155 ksi. As a result, thick walled tubes are often employed in certain formations in order to withstand the high collapse pressures present. However, the use of thick walled tubes significantly reduces the working space available for the explosive charges, which may limit the range of applications in which the tubes may be employed.

From the foregoing, then, there is a need for improved compositions for metallic tubular bars, and, in particular, systems and methods for producing metallic tubular bars with a combination of high tensile properties and toughness.

Embodiments of the invention are directed to steel tubes and methods of manufacturing the same. In one embodiment, a quenching and tempering procedure is performed in which a selected steel composition is formed and heat treated to yield a slightly tempered microstructure having a fine carbide distribution. In another embodiment, a double austenizing procedure is disclosed in which a selected steel composition is formed and subjected to heat treatment to refine the steel microstructure. In one embodiment, the heat treatment may comprise austenizing and quenching the formed steel composition a selected number of times (e.g., 2) prior to tempering. In another embodiment, the heat treatment may comprise subjecting the formed steel composition to austenizing, quenching, and tempering a selected number of times (e.g., 2). Steel products formed from embodiments of the steel composition in this manner (e.g., seamless tubular bars and pipes) will possess high yield strength, e.g., at least about 165 ksi, while maintaining good toughness.

In an embodiment, a steel tube is provided. The steel tube comprises

about 0.20 wt. % to about 0.30 wt. % carbon; about 0.30 wt. % to about 0.70 wt. % manganese; about 0.10 wt. % to about 0.30 wt. % silicon; about 0.90 wt. % to about 1.50 wt. % chromium; about 0.60 wt. % to about 1.00 wt. % molybdenum; about 0.020 wt. % to about 0.040 wt % niobium; and about 0.01 wt. % to about 0.04 wt. % aluminum; wherein the steel tube is processed to have a yield strength greater than about 165 ksi and wherein the Charpy V-notch energy is greater or equal to about 80 J/cm 2 in the longitudinal direction and greater than or equal to about 60 J/cm 2 in the transverse direction at about room temperature.

In a further embodiment, a method of making a steel tube is provided. The method comprises providing a carbon steel composition. The method further comprises forming the steel composition into a tube. The method also comprises heating the formed steel tube in a heating operation to a first temperature. The method additionally comprises quenching the formed steel tube in a quenching operation from the first temperature at a first rate such that the microstructure of the quenched steel is greater than or equal to about 95% martensite by volume. The method further comprises tempering the formed steel tube after the quenching operation by heating the formed steel tube to a second temperature less than about 550° C. The steel tube after tempering has a yield strength greater than about 165 ksi and the Charpy V-notch energy is greater or equal to about 80 J/cm 2 in the longitudinal direction and 60 J/cm 2 in the transverse direction at about room temperature.

In an additional embodiment, a method of forming a steel tube is provided. The method comprises providing a steel rod. The steel rod comprises

about 0.20 wt. % to about 0.30 wt. % carbon; about 0.30 wt. % to about 0.70 wt. % manganese; about 0.10 wt. % to about 0.30 wt. % silicon; about 0.90 wt. % to about 1.50 wt. % chromium; about 0.60 wt. % to about 1.00 wt. % molybdenum; about 0.020 wt. % to about 0.40 wt. % niobium; and about 0.01 wt. % to about 0.04 wt. % aluminum.

The method further comprises forming the steel rod into a tube in a hot forming operation at a temperature of about 1200° C. to 1300° C. The method further comprises heating the formed steel tube in a first heating operation to a temperature of about 880° C. to 950° C. for about 10 to 30 minutes. The method additionally comprises quenching the formed steel tube in a quenching operation after the first heating operation at a rate such that the microstructure of the quenched steel is greater than or equal to about 95% martensite. The method further comprises tempering the formed steel tube after the second quenching operation by heating the formed steel tube to a temperature between about 450° C. to about 550° C. for between about 5 minutes to about 30 minutes such that the final microstructure possesses about 95% martensite with the remainder consisting essentially of bainite. The microstructure, after tempering, may further include spherical carbides having a largest dimension less than or equal to about 150 μm and/or elongated carbides having a length less than or equal to about 1 μm and a thickness less than or equal to about 200 nm. The microstructure, after quenching, may further comprise an average grain size within the range between about 5 μm to about 15 μm. The steel tube after tempering has a yield strength greater than about 165 ksi and wherein the Charpy V-notch energy is greater or equal to about 80 J/cm 2 in the longitudinal direction and about 60 J/cm 2 in the transverse direction at about room temperature.

FIGS. 1A-1C are embodiments of methods of forming high strength steels;

FIGS. 2A-2B are micrographs of an embodiment of the steel composition after austenizing, quenching, and tempering heat treatments; and

FIG. 3 is a plot of Charpy impact energy (CVN) versus yield strength for steels formed from embodiments of the present disclosure.

Embodiments of the present disclosure provide steel compositions, tubular bars (e.g., pipes) formed using the steel compositions, and respective methods of manufacture. The tubular bars may be employed, for example, as perforating gun carriers for in the oil and gas industry. It may be understood, however, that tubular bars comprise one example of articles of manufacture which may be formed from embodiments of the steels 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 a generally hollow, elongate member which may be straight or have bends or curves and be formed to a predetermined shape, and any additional forming required to secure the formed tubular bar in its 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” refers to any elongate, hollow shape, which need not be circular or cylindrical.

The terms “approximately,” “about,” and “substantially” as used herein represent an amount close to the 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 carbon steels and methods of manufacture. In one embodiment, a selected steel composition is formed and subjected to heat treatment to refine the steel microstructure. In one embodiment, the steel composition may be formed and subjected to a heat treatment including austenizing, quenching, and tempering. The microstructure at the end of quenching includes at least about 95% martensite, by volume. Subsequent tempering may be performed within the range between about 450° C. to about 550° C. The microstructure resulting after tempering includes a fine carbide distribution, where the carbide particles are relatively small in size owing to the relatively low tempering temperatures. This microstructure provides relatively high strength and toughness. For example, yield strengths greater than about 165 ksi and Charpy V-Notch energies of at least 80 J/cm 2 in the LC direction and at least about 60 J/cm 2 in the CL direction.

In other embodiments, the heat treatment may comprise austenizing and quenching the formed steel composition a selected number of times (e.g., 2) to refine the grain size of the final microstructure. This refinement may improve the strength and toughness of the formed steel composition. Repeating the austenizing and quenching operations twice may be referred to herein as double austenizing. It may be understood, however, that the austenizing and quenching operations may be performed any number of times, without limit, to achieve the desired microstructure and mechanical properties. In another embodiment, the heat treatment may comprise subjecting the formed steel composition to austenizing, quenching, and tempering operations a selected number of times (e.g., 2), with tempering performed after each quenching operation.

It is anticipated that embodiments of articles formed from selected steel compositions in this manner (e.g., tubular bars and pipes) will possess high yield strength, at least about 165 ksi (about 1138 MPa), as measured according to ASTM E8, while maintaining good toughness. For example, experiments discussed herein illustrate that steels formed from embodiments of the disclosed composition may further exhibit Charpy V-notch impact energies greater than about 80 J/cm 2 in the LC direction and about 60 J/cm 2 in the CL direction, as measured according to ASTM Standard E23. As discussed in greater detail below, these improvements in properties are achieved, at least in part, due to refinement of the microstructure of the formed steel compositions (e.g., grain size, packet size, and average carbide size) as a result of varying the temperatures of respective austenizing operations.

For example, in one embodiment, repeated austenizing and quenching operations at different temperatures may be employed to refine the grain size and packet size of the formed steel tube with the objective of improving the toughness of the steel tube. The grain size of the tube can also be reduced by decreasing the austenizing temperature, as grain growth is a diffusion controlled process that may be delayed by reducing the austenizing temperature. However the austenizing temperature should also be high enough to decompose substantially all of the iron carbides (cementite) in the steel composition. If the austenizing temperature is not high enough, large cementite particles may remain in the final microstructure of the steel that impair the toughness of the steel. Thus, in order to improve the toughness of the steel, the austenizing temperature is preferably selected to be slightly above the minimum value to that is needed to dissolve the cementite. While temperatures higher than this minimum may guarantee the decomposition of cementite, they may produce excessive grain growth.

For this reason, a preferred temperature range for austenizing is provided in each condition. The preferred range depends on the iron carbide size of the initial microstructure. In an embodiment, if the steel is in the as hot-rolled condition (e.g., the case of the first austenizing treatment), the minimum temperature is preferably high enough to dissolve the large carbides appearing in the starting microstructure (e.g., about 900° C. to about 950° C.) If the material is in the as-quenched condition (e.g., the case of a second austenizing performed without intermediate tempering) there are substantially no cementite carbides present in the initial microstructure, so the minimum austenizing temperature is preferably lower (e.g., about 880° C. to about 930° C.).

These observations may be employed to reduce the austenizing temperature for refining the steel microstructure. If an intermediate tempering is performed, cementite carbides may be precipitated during tempering resulting in an increase in the minimum austenizing temperature as compared to the ideal case of the as quenched condition with substantially no cementite carbides.

However, during industrial processing it may be not possible or feasible to perform a double austenizing and quenching procedure without intermediate tempering. Therefore, the austenizing, quenching, and tempering operations may be repeated instead. When performing a tempering, reducing the tempering temperature is desirable in order to avoid the precipitation of large carbides, which need a higher austenizing temperature to be dissolved. For this reason, the tempering temperature is limited to less than about 550° C.

The metal composition of the present disclosure preferably comprises a steel alloy comprising not only carbon (C) but also manganese (Mn), silicon (Si), chromium (Cr), molybdenum (Mo), niobium (Nb), and aluminum (Al). Additionally, one or more of the following elements may be optionally present and/or added nickel (Ni), vanadium (V), titanium (Ti), and calcium (Ca). 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, sulfur (S), phosphorous (P), copper (Cu), nitrogen (N), lead (Pb), tin (Sn), arsenic (As), antimony (Sb), and bismuth (Bi). Elements within embodiments of the steel composition may be provided as below in Table 1, where the concentrations are in wt. % unless otherwise noted.

C is an element whose addition to the steel composition inexpensively raises the strength of the steel. In some embodiments, if the C content of the steel composition is less than about 0.20% it may be difficult to obtain the strength desired in the steel. On the other hand, in some embodiments, if the steel composition has a C content greater than about 0.30%, toughness may be impaired. Therefore, in an embodiment, the C content of the steel composition may vary within the range between about 0.20% to about 0.30%, preferably within the range between about 0.24% to about 0.27%.

Mn is an element whose addition to the steel composition is effective in increasing the hardenability, strength, and toughness. In some embodiments, if the Mn content of the steel composition is less than about 0.30%, it may be difficult to obtain the desired strength in the steel. However, in some embodiments, if the Mn content of the steel composition exceeds about 0.7%, banding structures within the steel may become marked and the toughness of the steel may decrease. Accordingly, in an embodiment, the Mn content of the steel composition may vary within the range between about 0.30% to about 0.7%, preferably within the range between about 0.45% to about 0.55%.

Si is an element whose addition to the steel composition has a deoxidizing effect during steel making process and also raises the strength of the steel. In some embodiments, if the Si content of the steel composition exceeds about 0.30%, the toughness and formability of the steel may decrease. Therefore, in an embodiment, the Si content of the steel composition may vary within the range between about 0.10% to about 0.30%, preferably within the range between about 0.20% to about 0.30%.

S is an impurity element whose presence within the steel composition causes the toughness and workability of the steel to decrease. Accordingly, in some embodiments, the S content of the steel composition is limited to less than or equal to about 0.010%, preferably less than or equal to about 0.003%.

P is an impurity element whose presence within the steel composition causes the toughness of the steel to decrease. Accordingly, in some embodiments, the P content of the steel composition limited to less than or equal to about 0.015%, preferably less than or equal to about 0.010%.

Cr is an element whose addition to the steel composition increases hardenability and tempering resistance of the steel. Therefore, Cr is desirable for achieving high strength levels. In an embodiment, if the Cr content of the steel composition is less than about 0.90%, it may be difficult to obtain the desired strength. In other embodiments, if the Cr content of the steel composition exceeds about 1.50%, the toughness of the steel may decrease. Therefore, in certain embodiments, the Cr content of the steel composition may vary within the range between about 0.90% to about 1.50%, preferably within the range between about 0.90% to about 1.0%.

Mo is an element whose addition to the steel composition is effective in increasing the strength of the steel and further assists in retarding softening during tempering. Mo additions to the steel composition may also reduce the segregation of phosphorous to grain boundaries, improving resistance to inter-granular fracture. In an embodiment, if the Mo content of the steel composition is less than about 0.60%, it may be difficult to obtain the desired strength in the steel. However, this ferroalloy is expensive, making it desirable to reduce the maximum Mo content within the steel composition. Therefore, in certain embodiments, Mo content within the steel composition may vary within the range between about 0.60% to about 1.00%, preferably within the range between about 0.65% to about 0.70%.

Ni is an element whose addition to the steel composition is optional and may increase the strength and toughness of the steel. However, Ni is very costly and, in certain embodiments, the Ni content of the steel composition is limited to less than or equal to about 0.50%, preferably less than or equal to about 0.15%.

Nb is an element whose addition to the steel composition may refine the austenitic grain size of the steel during hot rolling, with the subsequent increase in both strength and toughness. Nb may also precipitate during tempering, increasing the steel strength by particle dispersion hardening. In an embodiment, if the Nb content of the steel composition is less than about 0.020%, it may be difficult to obtain the desired combination of strength and toughness. However, in other embodiments, if the Nb content is greater than about 0.040%, a dense distribution of precipitates may form that may impair the toughness of the steel composition. Therefore, in an embodiment, the Nb content of the steel composition may vary within the range between about 0.020% to about 0.040%, preferably within the range between about 0.025% to about 0.030%.

V is an element whose addition to the steel composition may be used to increase the strength of the steel by carbide precipitations during tempering. However, in certain embodiments, V may be omitted from the steel composition. In an embodiment, when present, if the V content of the steel composition is greater than about 0.005%, a large volume fraction of vanadium carbide particles may be formed, with an attendant reduction in toughness of the steel. Therefore, in certain embodiments, the maximum V content of the steel composition may be less than or equal to about 0.005%.

Ti is an element whose addition to the steel composition may be used to refine austenitic grain size. However, in certain embodiments, Ti may be omitted from the steel composition. Additionally, in embodiments of the steel composition when Ti is present and in concentrations higher than about 0.010%, coarse TiN particles may be formed that impair toughness of the steel. Therefore, in certain embodiments, the maximum Ti content of the steel composition may be less than or equal to about 0.010%.

Cu is an impurity element that is not required in certain embodiments of the steel composition. However, depending upon the steel fabrication process, the presence of Cu may be unavoidable. Thus, in certain embodiments, the Cu content of the steel composition may be limited to less than or equal to about 0.30%, preferably less than or equal to about 0.15%.

Al is an element whose addition to the steel composition has a deoxidizing effect during the steel making process and further refines the grain size of the steel. In an embodiment, if the Al content of the steel composition is less than about 0.010%, the steel may be susceptible to oxidation, exhibiting high levels of inclusions. In other embodiments, if the Al content of the steel composition greater than about 0.040%, coarse precipitates may be formed that impair the toughness of the steel. Therefore, the Al content of the steel composition may vary within the range between about 0.010% to about 0.040%

Ca is an element whose addition to the steel composition is optional and may improve toughness by modifying the shape of sulfide inclusions. Thereafter, in certain embodiments, the minimum calcium content of the steel may satisfy the relationship Ca/S>1.5. In other embodiments of the steel composition, excessive Ca is unnecessary and the steel composition may comprise a Ca content less than or equal to about 0.05%.

The contents of unavoidable impurities including, but not limited to, S, P, N, Pb, Sn, As, Sb, Bi and the like are preferably kept as low as possible. However, mechanical properties (e.g., strength, toughness) of steels formed from embodiments of the steel compositions of the present disclosure may not be substantially impaired provided these impurities are maintained below selected levels. In one embodiment, the N content of the steel composition may be less than or equal to about 0.008%, preferably less than or equal to about 0.006%. In another embodiment, the Pb content of the steel composition may be less than or equal to about 0.005%. In a further embodiment, the Sn content of the steel composition may be less than or equal to about 0.02%. In an additional embodiment, the As content of the steel composition may be less than or equal to about 0.012%. In another embodiment, the Sb content of the steel composition may be less than or equal to about 0.008%. In a further embodiment, the Bi content of the steel composition may be less than or equal to about 0.003%.

In one embodiment, tubular bars may be formed using the steel composition disclosed above in Table 1. The tubular bars may preferably have a wall thickness selected within the range between about 4 mm to about 25 mm. In one embodiment, the metallic tubular bars may be seamless. In an alternative implementation, the metallic tubular bars may contain one or more seams.

Embodiments of methods 100 , 120 , 140 of producing high strength metallic tubular bars are illustrated in FIGS. 1A-1C . It may be understood that methods 100 , 120 , 140 may be modified to include greater or fewer steps than those illustrated in FIGS. 1A-1C without limit.

With reference to FIG. 1A , in operation 102 , the steel composition is formed and cast into a metallic billet. In operation 104 , the metallic billet may be hot formed into a tubular bar. In operations 106 (e.g., 106 A, 106 B, 106 C), the formed tubular bar may be subjected to heat treatment. In operation 110 , finishing operations may be performed on the bar.

Operation 102 of the method 100 preferably comprises fabrication of the metal and production of a solid metal billet capable of being pierced and rolled to form a metallic tubular bar. In one embodiment, the metal may comprise steel. In further embodiments, selected steel scrap 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, and 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 as understood in the art of steelmaking may be performed, and inclusion flotation may be 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 may produce 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 result. Following the production of the fluid slag, the steel may be cast into a round solid billet having a substantially uniform diameter along the steel axis.

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 1300° C., preferably about 1250° C. 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 may be 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 of about 1200° 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

In a non-limiting example, a solid bar possessing an outer diameter within the range between about 145 mm to about 390 mm may be hot formed as discussed above into a tube possessing an outer diameter within the range between about 39 mm to about 275 mm and a wall thickness within the range between about 4 mm to about 25 mm. The length of the tubes may be varied, as necessary. For example, in one embodiment, the length of the tubes may vary within the range between about 8 m to about 15 m.

In this fashion, a straight-sided, metallic tubular bar having a composition within the ranges illustrated in Table 1 may be provided.

In operations 106 A- 106 C, the formed metallic tubular bar may be subjected to heat treatment. In operation 106 A, a tubular bar formed as discussed above may be heated so as to substantially fully austenize the microstructure of the tubular bar. A tubular bar that is substantially fully austenized may comprise greater than about 99.9 wt. % austenite on the basis of the total weight of the tubular bar. The tubular bar may be heated to a maximum temperature selected within the range between about 880° C. to about 950° C. The heating rate during the first austenizing operation 106 A may vary within the range between about 15° C./min to about 60° C./min. The tubular bar may be further heated to the maximum temperature over a time within the range between about 10 minutes to about 30 minutes.

Following the hold period, the tubular bar may be subjected to quenching operation 106 B. In an embodiment, quenching may be performed using a system of water sprays (e.g., quenching heads). In another embodiment, quenching may be performed using an agitated water pool (e.g., tank) in which additional heat extraction is obtained by a water jet directed to the inner side of the pipe. In either case, the tubular bar may be cooled at a rate between approximately 15° C./sec to 50° C./sec to a temperature preferably not greater than about 150° C. The microstructure of the steel composition, after the quenching operation 104 , comprises at least about 95% martensite, with the remaining microstructure comprising substantially bainite.

Following the austenizing and quenching operations 106 A, 106 B, the tubular bar may be further subjected to a tempering operation 106 C. During the tempering operation 106 C, the tubular bar may be heated a temperature within the range between about 450° C. to about 550° C. The heating rate during the tempering operation 106 C may vary within the range between about 15° C./min to about 60° C./min. The tubular bar may be further heated to the maximum temperature over a time within the range between about 10 minutes to about 40 minutes. Upon achieving the selected maximum temperature, the tubular bar may be held at about this temperature for a time within the range between about 5 minutes to about 30 minutes.

Due to the low tempering temperatures, the final microstructure of the steel composition after the tempering operation 106 C comprises slightly tempered martensite having a fine carbide distribution. This microstructure is illustrated in FIGS. 2A-2B . As illustrated in FIG. 2 , the tempered martensite is composed of a ferrite matrix (e.g., dark gray phases) and several types of carbides (light gray particles).

With respect to morphology, two types of carbides were observed to be present in the microstructure, approximately spherical and elongated. Regarding the spherical carbides, the maximum size (e.g., largest dimension such as diameter) was observed to be about 150 nm. Regarding the elongated carbides, the maximum size was observed to be about 1 μm length and about 200 nm in thickness.

The hot rolled tube may be further subjected to different finishing operations 110 . Non-limiting examples of these operations may include cutting the tube to length, and cropping the ends of the tube, straightening the tube using rotary straightening equipment, if necessary, and non-destructive testing by a plurality of different techniques, such as electromagnetic testing or ultrasound testing. In an embodiment, the tubular bars may be straightened at a temperature not lower than the tempering temperature reduced by 50° C., and then cooled in air down to room temperature in a cooling bed.

Advantageously, seamless steel pipes obtained according to embodiments of the method 100 discussed above may be employed in applications including, but not limited to, perforating gun carriers in the oil and gas industry. As discussed in greater detail below, mechanical testing has established that embodiments of the steel pipes exhibit a yield strength of at least about 165 ksi (measured according to ASTM E8, “Standard Test Methods for Tension Testing of Metallic Materials,” the entirety of which is incorporated by reference) and a Charpy V-notch impact energy at room temperature, measured according to ASTM E23 (“Standard Test Methods for Notched Bar Impact Testing of Metallic Materials,” the entirety of which is incorporated by reference) of at least about 80 Joules/cm 2 for samples taken in the LC direction and at least about 60 Joules/cm 2 for samples taken in the CL direction.

The good combination of strength and toughness obtained in embodiments of the steel composition are ascribed, at least in part, to the combination of the steel composition and to the microstructure. In one aspect, the relatively small size of the carbides (e.g., spherical carbides less than or equal to about 150 nm and/or elongated carbides of about 1 μm or less in length and about 200 nm or less in thickness) increase the strength of the steel composition by particle dispersion hardening without strongly impairing toughness. In contrast, large carbides can easily nucleate cracks.

In alternative embodiments, one of methods 120 or 140 as illustrated in FIGS. 1B and 1C may be employed to fabricate seamless steel pipes when increased strength is desired. The methods 120 and 140 differ from one another and from the method 100 by the heat treatment operations performed on the seamless steel pipe. As discussed in greater detail below, embodiments of heat treatment operations 126 (of method 120 ) comprise repeated austenizing and quenching operations, followed by tempering. Embodiments of heat treatment operations 146 (of method 140 ) comprise repeated sequences of austenizing, quenching, and tempering. In other respects, the metal fabrication and casting, hot forming, and finishing operations of methods 100 , 120 , and 140 are substantially the same.

With reference to method 120 , the heat treatment 126 may comprise a first austenizing/quenching operation 126 A that may include heating and quenching a tubular bar formed as discussed above into the austenitic range. The conditions under which austenizing is performed during the first austenizing/quenching operation 126 A may be designated as A 1 . The conditions under which quenching is performed during the first austenizing/quenching operation 126 A may be designated as Q 1 .

In an embodiment, the first austenizing and quenching parameters A 1 and Q 1 are selected such that the microstructure of the tubular bar after undergoing the first austenizing/quenching operation 126 A comprises at least about 95% martensite with the remainder including substantially only bainite. In further embodiments, the first austenizing and quenching parameters A 1 and Q 1 may also produce a microstructure that is substantially free of carbides. In certain embodiments, a microstructure that is substantially free of carbides may comprise a total carbide concentration less than about 0.01 wt. % on the basis of the total weight of the tubular bar. In further embodiments, the average grain size of the tubular bar after the first austenizing and quenching operations 126 A may fall within the range between about 10 μm to about 30 μm.

In an embodiment, the first austenizing parameters A 1 may be selected so as to substantially fully austenize the microstructure of the tubular bar. A tubular bar that is substantially fully austenized may comprise greater than about 99.9 wt. % austenite on the basis of the total weight of the tubular bar. The tubular bar may be heated to a maximum temperature selected within the range between about 900° C. to about 950° C. The heating rate during the first austenizing operation 126 A may vary within the range between about 30° C./min to about 90° C./min. The tubular bar may be further heated to the maximum temperature over a time within the range between about 10 minutes to about 30 minutes.

The tubular bar may be subsequently held at the selected maximum temperature for a hold time selected within the range between about 10 minutes to about 30 minutes. The relatively low austenizing temperatures employed in embodiments of the disclosed heat treatments, within the range between about 900° C. to about 950° C., are employed to restrain grain growth as much as possible, promoting microstructural refinement that may give rise to improvements in toughness. For these austenizing temperatures, the austenizing temperature range of about 900° C. to about 950° C. is also sufficient to provide substantially complete dissolution of cementite carbides. Within this temperature range, complete dissolution of Nb- and Ti-rich carbides, even when using extremely large holding times, is generally not achieved. The cementite carbides, which are larger than Nb and Ti carbides, may impair toughness and reduce strength by retaining carbon.

Following the hold period, the tubular bar may be subjected to quenching. In an embodiment, quenching during the austenizing/quenching operations 126 A may be performed a system of water sprays (e.g., quenching heads). In another embodiment, quenching may be performed using an agitated water pool (e.g., tank) in which additional heat extraction is obtained by a water jet directed to the inner side of the pipe.

Embodiments of the quenching parameters Q 1 are as follows. The tubular bar may be cooled at a rate between approximately 15° C./sec to 50° C./sec to a temperature preferably not greater than about 150° C.

The second austenizing/quenching operation 126 B may comprise heating and quenching the tubular bar formed as discussed above into the austenitic range. The conditions of under which austenizing is performed during the second austenizing/quenching operation 126 B may be designated as A 2 . The conditions under which quenching is performed during the second austenizing/quenching operation 126 B may be designated as Q 2 .

In an embodiment, the second austenizing and quenching parameters A 2 and Q 2 may be selected such that the microstructure of the tubular bar after undergoing the second austenizing/quenching operation 126 B comprises at least about 95% martensite. In further embodiments, the austenizing and quenching parameters A 2 and Q 2 may also produce a microstructure that is substantially free of carbides.

In additional embodiments, the average grain size of the tubular bar after the second austenizing/quenching operations 126 B may be less than that obtained after the first austenizing and quenching operations 126 A. For example, the grain size of the tubular pipe after the second austenizing/quenching operations 126 B may fall within the range between about 5 μm to about 15 μM. This microstructural refinement may improve the strength and/or the toughness of the tubular bar.

In an embodiment, the second austenizing parameters A 2 are as follows. The tubular bar may be heated to a maximum austenizing temperature less than that employed in the first austenizing/quenching operations 126 A in order to further refine the grain size of the microstructure. The second austenizing operation A 2 takes advantage of the carbide dissolution achieved during the first austenizing/quenching operations 106 A (A 1 /Q 1 ). As substantially all the iron carbides (e.g., cementite particles) are dissolved within the microstructure following the first austenizing and quenching operations 126 , lower austenizing temperatures can be used during the second austenizing and quenching operations 126 B with attendant reduction in grain size (grain refinement). In an embodiment, the second austenizing operation A 2 may take place at a temperature selected within the range between about 880° C. to about 930° C. The heating rate during the second austenizing operation A 2 may vary within the range between about 15° C./min to about 60° C./min. The tubular bar may be subsequently held at the selected maximum temperature for a hold time selected within the range between about 10 to about 30 minutes.

Following the hold period, the tubular bar may be subjected to quenching Q 2 . In an embodiment, quenching during the austenizing/quenching operations 126 B may be performed a system of water sprays (e.g., quenching heads). In another embodiment, quenching may be performed using an agitated water pool (e.g., tank) in which additional heat extraction is obtained by a water jet directed to the inner side of the pipe.

Embodiments of the quenching parameters Q 2 are as follows. The tubular bar may be cooled at a rate between about 15° C./sec to about 50° C./sec to a temperature preferably not greater than about 150° C.

Following the first and second austenizing/quenching operations 126 A, 126 B, the tubular bar may be further subjected to a tempering operation 126 C, also referred to herein as (T). During the tempering operation 126 C, the tubular bar may be heated a temperature within the range between about 450° C. to about 550° C. The heating rate during the tempering operation 106 C may vary within the range between about 15° C./min to about 60° C./min. The tubular bar may be further heated to the maximum temperature over a time within the range between about 10 minutes to about 40 minutes. Upon achieving the selected maximum temperature, the tubular bar may be held at about this temperature for a time within the range between about 5 minutes to about 30 minutes.

The tubular bars may also be subjected to finishing operations 130 . Examples of finishing operations 130 may include, but are not limited to, straightening. Straightening may be performed at a temperature not lower than the tempering temperature reduced by 50° C. Subsequently the straightened tube may be cooled in air down to about room temperature in a cooling bed.

In an alternative embodiment, the formed tubular bar may be subjected to method 140 which employs heat treatment operations 146 C. In heat treatment operations 146 C, first austenizing and quenching operations 146 A (A 1 ) and (Q 1 ) are followed by a first tempering operation 146 B (T 1 ), second austenizing and quenching operations 146 C (A 2 ) and (Q 2 ), and second tempering operation 146 D (T 2 ). The first and second austenizing and quenching operations 146 A and 146 C may be performed as discussed above with respect to the first and second austenizing and quenching operations 126 A and 126 B. The first (T 1 ) and second (T 2 ) tempering operations 146 B and 146 D may also be performed as discussed above with respect to the first tempering operation 106 C.

The microstructure resulting from methods 120 and 140 may be similar to that resulting from method 100 . For example, in one embodiment, after the first austenizing and quenching operations 126 A and 146 A, the average grain size may vary within the range between about 10 μm to about 30 μm. In another embodiment, after the second austenizing and quenching operations 126 C and 146 C, the average grain size may vary within the range between about 5 μm to about 15 μm. In further embodiments, a fine distribution of carbides may be present within the microstructure after tempering operations 126 C, 146 D. For example, spherical and elongated carbides may be present within the microstructure, with the maximum size of the spherical particles being less than or equal to about 150 nm and the maximum size of the elongated carbides being less than or equal to about 1 μm length and less than or equal to about 200 nm in thickness.

Advantageously, seamless steel pipes and tubes formed according to the embodiments of methods 120 and 140 may be suitable for applications including, but not limited to, perforating gun carriers in the oil and gas industry. For example, in one embodiment, tubular bars and pipes formed from embodiments of the steel composition may exhibit a yield strength of at least about 170 ksi (about 1172 MPa) as measured according to ASTM Standard E8. In another embodiment, tubular bars and pipes formed from embodiments of the steel composition may exhibit Charpy V-notch impact energies at room temperature greater than about 80 J/cm 2 in the LC direction and about 60 J/cm 2 in the CL direction as measured according to ASTM Standard E23. This good combination of properties is ascribed, at least in part, to the refined grain size and relatively small size of the carbides within the microstructure.

Beneficially, in certain embodiments, these results may be achieved without vanadium addition. Vanadium is known to increase strength by carbide precipitation during tempering but may impair toughness.

In the following examples, the tensile and impact properties of steel pipes formed using embodiments of the steel making method discussed above are illustrated. The formed steel pipes were tested after heat treatments of austenizing, quenching, and tempering (A+Q+T) (Conditions 1 and 2 ), double austenizing and tempering (A 1 +Q 1 +A 2 +Q 2 +T) followed by tempering (Condition 3 ). The tested steel pipes possessed an outer diameter of about 114.3 mm and a wall thickness of about 8.31 mm, unless otherwise noted. Experiments were performed on samples having approximately the composition and heat treatments of Tables 2 and 3, respectively.

Measurements of strength and impact properties were performed on between 3 to 5 pipes for each condition. For each tube, tensile tests were performed in duplicate and impact tests were performed in triplicate at about room temperature. It may be understood that the examples presented below are for illustrative purposes and are not intended to limit the scope of the present disclosure.

Room Temperature Tensile Properties and Impact Energies

The strength and elongation of steels having compositions as indicated above in Tables 2 and 3 at were measured according to ASTM Standard E8 at room temperature. The Charpy energies of the steels of Tables 2 and 3 were measured according to ASTM Standard E23 at about room temperature and represent a measure of the toughness of the materials. The Charpy tests were performed on samples having dimensions of about 10×7.5×55 mm taken longitudinally (LC) from the pipes. The average tensile strength, yield strength, elongation, and Charpy V-notch energies (CVN) measured for each condition are reported in Table 4 and average values per tube are reported in FIG. 3 .

For each of the conditions tested, yield strength was observed to be greater than or equal to about 165 ksi and ultimate tensile strength was observed to be greater than or equal to about 170 ksi. The elongation at failure for each of the conditions tested was further found to be greater than or equal to about 10%. In further embodiments, the yield strength was observed to be greater than about 170 ksi, ultimate tensile strength was observed to be greater than or equal to about 180 ksi, and elongation at failure was found to be greater than or equal to about 13%. In certain embodiments, the measured Charpy V-notch impact energies at about room temperature were greater than about 65 J/cm 2 for each of the conditions tested. In further embodiments, the room temperature Charpy energies were greater than or equal to about 90 J/cm 2 .

The best combination of tensile properties and toughness were observed for heat treatment condition 3 , which corresponded to double austenizing. This condition exhibited the largest yield strength (about 189 ksi) and CVN at room temperature (about 97 J/cm 2 ). The improvement in yield strength and toughness is ascribed to the microstructural refinement achieved by the double austenizing/quenching operations.

Further Impact Energy Studies

Additional impact energy investigations were performed on steel pipe samples formed according to Condition 1 from about −60° C. to about room temperature in order to identify the ductile to brittle transition temperature of the formed steel compositions. For these measurements, samples were taken in both the longitudinal (LC) and transverse (CL) directions, Charpy tests were performed on samples having dimensions of about 10×7.5×55 mm in the LC orientation and about 10×5×55 mm in the CL orientation. The average Charpy V-notch energies for each condition are reported in Table 5.

As illustrated in Table 5, the LC Charpy samples at about room temperature (RT) exhibited energies greater than about 80 J/cm 2 and approximately 100% ductile fracture, as observed from the fracture surface. The CL Charpy samples exhibited energies of greater than about 60 J/cm 2 and approximately 100% ductile fracture. As the test temperature decreased from about room temperature to about −60° C., the LC and CL Charpy energies dropped by roughly half to approximately 30-36 J/cm 2 . Concurrently, the portion of the fracture surface undergoing ductile fracture decreased by approximately two-thirds in each geometry.

From the results, it can be observed that the ductile to brittle transformation temperature (DBTT) is between −20° C. and −40° C. for longitudinally oriented samples (LC) owing to the large reduction in ductile area observed between about −20° C. and about −40° C. in the LC orientation (from about 71% to about 44%). It can be further observed that the DBTT is between about −40° C. and −60° C. for transversely oriented samples (CL) owing to the large reduction in ductile area observed between about −40° C. and about −60° C. (from about 75% to about 31%).

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.