EP2287346B1

EP2287346B1 - Bainitic steels with boron

Info

Publication number: EP2287346B1
Application number: EP10166261.7A
Authority: EP
Inventors: Perez Teresa Estela; Gomez Gonzalo Roberto
Original assignee: Tenaris Connections BV
Current assignee: Tenaris Connections BV
Priority date: 2009-06-17
Filing date: 2010-06-17
Publication date: 2019-12-18
Anticipated expiration: 2030-06-17
Also published as: JP2011006790A; US20100319814A1; EP2287346A1; MX2010006761A; BRPI1004267A2; AR077129A1; BRPI1004267B1; JP5787492B2

Description

BACKGROUND OF THE INVENTION

Field

Embodiments of the present disclosure pertain to seamless pipes formed from steels containing micro-alloying additions of boron and titanium, with yield strengths of at least 100 ksi (690 MPa), excellent toughness, and good weldability. Such pipes are suitable for use as high strength line pipes, for example X100 in API 5L standard, and other possible applications.

Description of the Related Art

Micro-alloying additions of boron to steel are desirable, as such additions may improve the mechanical properties of the steel. For example, boron additions may increase hardenability, the ability of steel to be hardened by heat treatment. By migrating to grain boundaries, where they inhibit austenite to ferrite phase transformation, boron additions may improve the ease with which martensite may be formed. Furthermore, boron is effective at very low concentrations, providing significant improvements in hardenability at relatively low cost.
In order to achieve these benefits, boron should remain in its free, elemental state. However, boron reacts easily with impurities present in the steel, such as nitrogen. When boron nitrides are formed, the positive effect on hardenability provided by boron may be reduced, owing to the decrease in free boron.
To address this issue, strong nitride formers, such as titanium, may be added to the steel composition in order to inhibit boron nitrides from forming. Concomitantly, however, relatively coarse titanium nitride particles may be formed during solidification. These particles, which may further grow during reheating prior to hot rolling, can lead to poor toughness in the steel and overshadow the property improvements yielded by the boron addition. EP2028284 relates to a high-strength seemless steel pipe for a mechanical structure which has excellent toughness and weldability, and method for manufacture thereof.

SUMMARY

A method of making a steel pipe is provided according to claim 1.
In certain embodiments, the steel composition may be formed into a seamless pipe.
In a further embodiment, a method of making a steel pipe is provided according to claim 5.
In certain embodiments, the microstructure may comprise no more than about 5% of martensite.
In a further embodiment, a method of making a steel pipe is provided according to claim 11.
In certain embodiments, the air cooled composition may comprise no less than about 50% of martensite.

BRIEF DESCRIPTION OF THE DRAWINGS

Figure 1 is a schematic flow diagram of one embodiment of a method of producing boron-titanium (B/Ti) steel pipes;
Figure 2 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 1;
Figure 3 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 1 cooled from the austenitic range at rates of about 2°C/sec, 5°C/sec, 10°C/sec, and 20°C/sec;
Figures 4A and 4B are plots of impact energy (CVN) for an embodiment of steel composition 1 subjected to accelerated cooling; (A) impact energy as a function of cooling rate; (B) impact energy as a function of temperature;
Figure 5 is a plot of hardness as a function of tempering temperature for an embodiment of steel composition 1 in the quenched and tempered condition;
Figure 6 illustrates a scanning electron micrograph of the microstructure of an embodiment of steel composition 1 that is quenched and tempered at about 410°C;
Figure 7 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 2;
Figure 8 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 2 cooled from the austenitic range at rates of about 0.2°C/sec, 0.5°C/sec, 1°C/sec, 10°C/sec, 30°C/sec, and 50°C/sec;
Figure 9 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 3;
Figure 10 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 3 cooled from the austenitic range at rates of about 0.2°C/sec, 0.5°C/sec, 1°C/sec, 10°C/sec, 30°C/sec, and 50°C/sec;
Figures 11A-11B are plots of hardness as a function of cooling rate from hot rolling for embodiments of steel compositions 2 and 3; (A) composition 2; (B) composition 3;
Figures 12A-12B illustrate scanning electron micrographs of the microstructure of embodiments of steel compositions 2 and 3 in the as-quenched condition; (A) composition 2; (B) composition 3;
Figures 13A-13B illustrate scanning electron micrographs of the microstructure of embodiments of steel compositions 2 and 3 in the quenched and tempered condition; (A) composition 2; (B) composition 3;
Figure 14 is a plot of hardness as a function of tempering temperature for embodiments of steel compositions 2 (solid squares) and 3 (open squares); and
Figure 15 is a plot of hardness as a function of the average cooling rate between 800°C and 500°C for an embodiment of steel composition 2 steel and a reference Nb-V steel.

DETAILED DESCRIPTION

Embodiments of the present disclosure present compositions and methods of manufacture for low carbon steels microalloyed with boron. In particular, boron/titanium (B/Ti) steels which exhibit controlled particulates of titanium nitride (TiN), and attendant improvements in toughness, are discussed in detail. Through addition of titanium and boron, free boron may be substantially kept in solid solution, improving hardenability during austenite decomposition.
The size of TiN precipitates may be controlled by the cooling rate during casting. In certain embodiments, size may comprise the diameter of the precipitates. In other embodiments, size may comprise the largest dimension of the precipitates. For example, as discussed in detail below, by employing cooling rates greater than about 10 to 30°C/min during casting, fine precipitates of TiN, having a mean size less than about 50 nm, may be produced. Due to the small size of these TiN precipitates, they are not detrimental to toughness. Additionally, these precipitates may inhibit excessive grain growth during processing operations such as reheating prior to hot rolling. By reducing austenite grain size, toughness may be improved, after accelerated cooling or quenching, due to the reduction in martensite/bainite packet size.
The mechanical properties and microstructure of the steel composition may be further influenced by heat treatments after hot rolling. In one embodiment, steel compositions may be cooled in air at rates less than about 1°C/sec after hot rolling and subjected to reheating into the austenitic range and quenching. In other embodiments, steel compositions may be cooled in air after hot rolling and subjected to reheating into the austenitic range and quenching and tempering. In further embodiments, steel compositions may be subject to accelerated cooling at rates between about 5 to 50°C/sec directly after hot rolling.
Excellent combinations of mechanical properties may be obtained for compositions processed in this manner, especially in the case of compositions subjected to quenching and tempering. For example, samples subjected to quenching and tempering at about 500°C may exhibit yield and tensile strengths of about 118 and 127 ksi, respectively, with impact energies measured in the range of about 143-173 J at about -60°C.
In another example, samples subjected to accelerated cooling may exhibit good impact energies, especially for cooling rates of about 10-20°C/sec. For example, impact energies greater than about 220 J are observed for temperatures of -20°C and higher. These and other advantages of the disclosed embodiments are discussed in detail below.
Figure 1 illustrates one embodiment of a method 100 of producing boron-titanium (B/Ti) steels. In certain embodiments, the compositions may be produced in the form of pipes. The method 100 of Figure 1 includes steel casting operations in blocks 110, 112, and 114, collectively referred to as steel casting operations 102, steel forming operations in blocks 116, 120, 122 and 124, collectively referred to as steel forming operations 104, and steel heat treatment operations in blocks 126 and 128, collectively referred to as heat treatment operations 106. It may be appreciated that, in some embodiments one or more of the heat treatment operations can be omitted partially or totally, as necessary.
The B/Ti steel is cast from the molten state during steel casting operations 102. In certain embodiments, the steel casting operations 102 may comprise continuous casting operations. For example, the steel casting operations 102 can include iron melting/purification 110, ladle treatments 112, and continuous casting 114, as are known in the art.

In one embodiment, the steel may comprise elements in the concentration ranges listed below in Table 1, where the concentrations are provided in weight percent (wt. %) on the basis of the total weight of the steel composition, unless otherwise noted.

Table 1 - Steel Composition

Element	Concentration (wt. %)
Element	Minimum	Maximum
C	0.04	0.12
Mn	0.6	1.6
Si	0.05	0.3
Ti	0.01	0.03
B	0.0005	0.003
Ni	0	0.5
Cr	0	0.5
Mo	0	0.5
V	0	0.15
Nb	0	0.05
N	0	0.008
Ratio Ti/N	> 3.4

The concentration of the elements may be further selected such that the carbon equivalency, CE_Pcm, of the composition is less than about 0.22, where CE_Pcm is calculated according to:

CE = C + \frac{Si}{30} + \frac{Mn + Cu + Cr}{30} + \frac{Ni}{60} + \frac{Mo}{15} + \frac{V}{10} + 5 B

where the concentration of each element is provided in wt. %.

As illustrated in Table 1, the cast steel may comprise a boron-titanium steel alloy including not only carbon (C), boron (B), and titanium (Ti) but one or more of manganese (Mn), silicon (Si), nickel (Ni), chromium (Cr) molybdenum (Mo), vanadium (V), and niobium (Nb). Impurities of sulfur (S), phosphorous (P), copper (Cu), and nitrogen (N) may also be present, however, the concentration of these impurities in one embodiment is preferably reduced to an amount as low as possible.
C is an element whose addition inexpensively raises the strength of the steel. If the C content is less than about 0.04%, it may be, in some embodiments, difficult to obtain the strength desired in the composition. On the other hand, in other embodiments, if the steel has a C content greater than about 0.12 wt. %, toughness and weldability may be adversely impacted. Therefore, in an embodiment, the C content may range between about 0.04 to 0.12 wt. %. In other embodiments, the C content may range between about 0.04 to 0.08 wt. %. This lower C range may enable compositions to be fabricated, optionally, without tempering (i.e. in the as-quenched condition), while still achieving good toughness.
B is an element whose addition is effective in increasing the hardenability of the steel. For example, B may improve hardenability by inhibiting the formation of ferrite. If the B content is less than about 0.0005 wt. %, in some embodiments, it may be difficult to obtain the desired hardenability of the steel. However, if the B content too high, in other embodiments, coarse boron carbides may be formed at grain boundaries, adversely affecting toughness. Accordingly, in an embodiment, the concentration of B in the composition may range between about 0.0005 to 0.003 wt. %. In other embodiments, the concentration of B in the composition may range between about 0.0005 to 0.002 wt. %. At least a portion of the B in the composition may be in its free, elemental state in solid solution.
Si is an element whose addition has a deoxidizing effect during the steel making process and also raises the strength of the steel. If the Si content is too low, in some embodiments, the steel may be susceptible to oxidation, with a high level of micro-inclusions. On the other hand, though, if the Si content of the steel is too high, in some embodiments both toughness and formability of the steel may decrease. Therefore, in certain embodiments of the composition, the concentration of Si may range between about 0.05 to 0.3 wt. %.
Mn and Cr are elements which may be employed in combination with B, Mo, and Ni to increase hardenability. For example, these alloying additions may assist in inhibiting the formation of ferrite and pearlite from austenite during cooling. They may further enable depression of the bainitic start temperature, improving microstructural refinement. Mn may additionally provide solid solution hardening. In certain embodiments, the concentration of Mn may range between about 0.6 to 1.6 wt. %. In further embodiments, Cr may be omitted from the composition. In other embodiments, the concentration of Cr may range up to about 0.5 wt. %.
Mo is an element used to increase the hardenability of the steel composition. Alloying additions of Mo may also reduce the segregation of phosphorous to grain boundaries, improving resistance to inter-granular fracture. Mo may further enhance the hardenability effects of B. In certain embodiments, Mo may be omitted from the composition. In other embodiments, the concentration of Mo may range up to about 0.5 wt. %.
Ni is an alloying addition which may increase hardenability and improve toughness. In certain embodiments, Ni may be omitted from the composition. In other embodiments, the concentration of Ni may range up to about 0.5 wt. %.
Ti is an element whose addition is effective in increasing the effectiveness of B in the steel, by fixing nitrogen impurities as TiN and inhibiting the formation of boron nitrides. If the Ti content is too low it may be difficult, in some embodiments, to obtain the desired effect of boron on hardenability. In an embodiment, if the Ti content is higher than about 0.03 wt. %, coarse TiN and TiC may be formed, adversely affecting hot ductility and toughness. Accordingly, in certain embodiments, the concentration of Ti may range between about 0.01 to 0.03 wt. %.
In alternative embodiments, the concentration of Ti may be specified on the basis of the concentration of N, maintaining a ratio of Ti to N greater than about 3.4 (for concentrations in weight percent).
In certain embodiments, substantially all of the N present within the composition may be in the form of TiN. In certain embodiments, greater than about 90%, greater than about 92%, greater than about 94%, greater than about 96%, greater than about 98%, and greater than about 99% of the N content of the composition may be present in the form of TiN. The TiN may adopt forms including, but not limited to, particles.
Nb is an alloying addition which may be used to refine the austenitic grain size of the composition. Nb may further enhance the effects of boron on hardenability and provide precipitation hardening. In certain embodiments, Nb may be omitted from the composition. In other embodiments, the concentration of Nb may range up to about 0.05 wt. %.
V is an alloying addition that may be employed to provide precipitation hardening. In certain embodiments, V may be omitted from the composition. In other embodiments, the concentration of V may range up to about 0.15 wt. %.
O is an impurity which may be present in the steel composition, for example, in the form of oxides. As the oxygen content increases, impact properties may be impaired. Accordingly, a lower oxygen content is preferred. In one embodiment, the upper limit of the oxygen content may be about 0.0050 wt. %. In another embodiment, the upper limit of oxygen content is below about 0.0015 wt. %.
Cu is not needed in embodiments of the steel composition, but may be present. In some embodiments, depending on the manufacturing process, the presence of Cu may be unavoidable. Thereafter, in an embodiment, the maximum Cu content may be about 0.10 wt. % or less.
S, P, Ca, N, and the like are impurities and their concentration is preferably kept as low as possible. In certain embodiments, the concentration of each of S, P, Ca, and N may be independently provided as: S not greater than about 0.005 wt. %, P not greater than about 0.015 wt. %, Ca not greater than about 0.003 wt. %, and N not greater than about 0.008 wt. %. In alternative embodiments the concentration of each of S, P, Ca, and N may be independently provided as: S not greater than about 0.003 wt. %, P not greater than about 0.015 wt. %, Ca not greater than about 0.002, and N not greater than about 0.006 wt. %.
The liquid steel may be continuously cast in steel casting operation 114. In certain embodiments, the liquid steel may be cast into a rod, although, it may be understood that other shapes may be cast. In particular, the cooling rate of the cast rod may be selected so as to provide control over the size of TiN precipitates that form during solidification. In certain embodiments, in order to inhibit coarsening of the TiN precipitates, the cooling rate during casting may be maintained at a selected rate. The cooling rate is selected such that the size of the TiN precipitates is less than about 50 nm. In an embodiment, the cooling rate from casting may be maintained at a rate greater than about 5°C/min at about the center of the rod. In further embodiments, the cooling rate from casting may be maintained at a rate greater than about 10°C/min at about the center of the rod. In other embodiments, the cooling rate from casting may be maintained at a rate greater than about 20°C/min at about the center of the rod. In additional embodiments, the cooling rate from casting may be maintained at a rate greater than about 30°C/min at about the center of the rod.
In one embodiment, the rod thus fabricated may be subsequently formed into a tubular bar or pipe in steel forming operations 104, and more particularly may be formed into a seamless pipe. A solid, substantially cylindrical rod of steel may be subjected to a first reheating operation (block 116) into the austenitic range, up to a temperature of about 1200°C to 1300°C, preferably about 1250°C. In blocks 120 and 122, the rod may be further pierced, in certain preferred embodiments, utilizing the Mannessmann process at temperatures between about 1100 to 1200°C, and subsequently hot rolled at temperatures ranging between about 900 to 1100°C.
Advantageously, the seamless hot rolled tube of steel may possess an approximately uniform wall thickness, both circumferentially around the tube and longitudinally along the tube axis. In one example, tubes formed in this manner may possess an outer diameter ranging between about 60 to 273 mm and wall thickness ranging between about 6 to 25 mm. In another example, a solid bar possessing an outer diameter of about 290 mm may be hot rolled in this manner into a tube possessing an outer diameter of about 244.5 mm and a wall thickness of about 16 mm.
During hot rolling, the cross-sectional area reduction experienced by the tube may provide a refined microstructure. A refined microstructure advantageously allows obtaining desired mechanical properties within the fabricated tube. The seamless hot rolled tube of steel so manufactured may then be cooled to room temperature. In certain embodiments, the austenitic grain size of the steel, after hot rolling and prior to transformation, may range between about 10 to 50 µm. In other embodiments, the austenitic grain size of the steel, after hot rolling and prior to transformation, may range between about 20 to 50 µm.
Beneficially, this degree of austenitic refinement may allow selected compositions to achieve a good balance of strength and toughness after accelerated cooling from the finish rolling temperature without the need for subsequent heat treatment, such as quenching or quenching and tempering. In other embodiments, this degree of austenitic refinement may allow compositions having elevated carbon concentrations to achieve a good balance of strength and toughness when subjected to heat treatments such as quenching and tempering.
Embodiments of the composition may be cooled from hot rolling by air cooling or accelerated cooling in block 124. When cooling from air, cooling rates less than about 1°C/sec may be achieved in tubes with wall thickness greater than about 8 mm. Subsequent heat treatments may also be employed to improve the strength and toughness of the steel composition.
When performing accelerated cooling, cooling may be performed directly from hot rolling, without an intermediate cooling step, to room temperature (block 124). Several devices can be used to achieve cooling rates greater than that corresponding to natural air cooling, including, but not limited to, forced air flow, water sprays, and air-water mixture sprays. The flow of the coolant may be directed to the outer tube wall, or to the inner and outer tube walls in order to improve microstructure homogeneity. In certain embodiments, using the above mentioned cooling alternatives, cooling rates ranging between about 5 to 50°C/sec may be achieved through accelerated cooling. In other embodiments, cooling rates ranging between about 10 to 50°C/sec may be employed. In further embodiments, cooling rates ranging between about 10 to 20°C/sec may be employed. In additional embodiments, these cooling rates may be employed with tubes of wall thickness between about 8 mm and 25 mm.
Forming of the hot rolled tube may be completed through a plurality of finishing steps. Non-limiting examples of the finishing steps may include cutting the tube to length, such as lengths of approximately 8 m to 15 m, cropping the ends of the tube, straightening the tube, and non-destructive testing (e.g., electromagnetic testing, ultrasound testing). In this fashion, a substantially straight-sided, metallic tubular bar having a composition within the ranges illustrated in Table 1 may be provided.
One or more heat treatment operations 106 may optionally be performed upon the tube after the forming operations 104. In an embodiment, quenching may be performed in block 126. For example, the composition may be reheated a second time into the austenitic range, prior to quenching, to temperatures greater than about Ae3 (e.g., about 870-950°C). Soak times at maximum temperature may range between about 5 to 30 minutes. Quenching may be further performed with water sprays to cool the composition from about the maximum temperature to about room temperature.
In other embodiments, tempering may be further performed upon quenched compositions in block 128. Tempering may be performed by heating to temperatures ranging between about 400 to 700°C, holding at the tempering temperature for a selected duration, and air cooling from the tempering temperature to about room temperature. The compositions may be held at the tempering temperature for between about 10 to 60 minutes.

Examples

The manufacture, microstructure, and mechanical properties of embodiments of three steel compositions of the present disclosure, referred to as compositions 1, 2, and 3, are discussed in the examples below. The performance benefits achieved from such compositions are further discussed. It may be understood that these examples are discussed for illustrative purposes and should not be construed to limit the scope of the disclosed embodiments.

The concentrations of alloying elements present in

compositions

1, 2, and 3 are illustrated below in Table 2.

Table 2 - Composition of steel compositions 1, 2, and 3

	Composition 1	Composition 2	Composition 3
C (wt%)	0.07	0.09	0.06
Mn (wt%)	1.41	1.12	1.18
Si (wt%)	0.26	0.08	0.09
Ni (wt%)	0.37	0.02	0.02
Cr (wt%)	0.29	0.02	0.31
Mo (wt%)	0.28	0.26	0.27
V (wt%)	0.062	0.002	0.003
Nb (wt%)	0.031	0.003	0.037
Ti (wt%)	0.018	0.024	0.023
B (ppm)	10	13	18
CE(P_cm)	0.20	0.174	0.165

CE (P_{cm}) = [CE (P_{cm}) = C + Si / 30 + (Mn + Cu + Cr) / 20 + Ni / 60 + Mo / 15 + V / 10 + 5 B]

Composition 1 was designed to produce a fine bainitic structure after accelerated cooling from the austenitic range. In contrast, composition 2, having the highest carbon content, was designed for use with air cooling from hot rolling, followed by quenching and tempering, as discussed below. Due to its higher carbon content, Cr and Nb are substantially absent in composition 2, as compared with the other compositions. Further, composition 3, having the lowest carbon content, was designed to obtain high strength and good toughness in the as-quenched condition, without tempering.
Steels having compositions 1, 2 and 3 were melted in an approximately 20 kg vacuum induction furnace and electro-slag re-melted to decrease sulfur content. Subsequently, compositions 1, 2, and 3 were cast into slabs having a thickness of approximately 140 mm and hot rolled to a final thickness of about 16 mm. During hot rolling, reheating and finishing temperatures of about 1200-1250°C and 950-1000°C, respectively, were employed. All the hot-rolled plates were subsequently air cooled to about room temperature.
The hot rolled composition 1 was subjected to one of the following post-rolling processing sequences:

a) reheating into the austenitic range at temperatures about 900-950°C, followed by quenching;
b) reheating into the austenitic range at temperatures about 900-950°C, followed by quenching and tempering;
c) reheating into the austenitic range at temperatures about 920-950°C, followed by accelerated cooling.

Quenching was performed in water, from a temperature of about 900-950°C to about room temperature, using moderate agitation. Where tempering operations were also performed, the composition was heated to between about 300°C to 450°C with a soaking time of about one hour.
Accelerated cooling was performed by cooling the composition in a mixture of air and water from the reheating temperature of about 900-950°C, to about room temperature, at cooling rates ranging between about 5 to 45°C/sec. The reheating temperature prior to accelerated cooling was selected to have an austenitic microstructure representative of that industrially achieved just at the exit of the hot rolling mill. The complete heat treatment was performed using a Gleeble 3500 thermo-mechanical simulator.
Hot rolled compositions 2 and 3 were subject to one of two post-rolling processing sequences:

a) reheating into the austenitic range followed by quenching.
b) reheating into the austenitic range followed by quenching and tempering.

Quenching was performed by heating the composition to a temperature of about 925°C (composition 2) or about 930°C (composition 3), with a soak time of about 10 min. Cooling was performed in water from the quenching temperature to about room temperature using moderate agitation. When tempering was performed, the composition was heated to temperatures ranging between about 400 to 700°C, with a soak time at maximum temperature of about 30 minutes.
For each composition, dilatometric tests were performed using a Gleeble 3500 thermomechanical simulator to evaluate the continuous cooling transformation (CCT) behavior. Compositions 1, 2, and 3 were reheated at about 5°C/sec to about 920°C, 925°C, and 930°C, respectively, with a soak time of about 10 min at maximum temperature. The austenization temperatures were chosen to be approximately 20-30°C above the Ac3 temperature corresponding to the respective compositions. Cooling rates ranging between about 0.5 to 50°C/sec were examined in composition 1 and cooling rates ranging between about 0.2 to 50°C/sec were examined in compositions 2 and 3. The resulting microstructures were further characterized using optical and scanning electron microscopy.
The mechanical properties of the compositions so fabricated were further evaluated by mechanical tests including one or more of tensile testing, hardness testing, and Charpy testing. In each case, tensile samples and full size Charpy samples were taken in the transversal direction. Tensile testing was performed in accordance with ASTM E8, "Standard Test Methods for Tension Testing of Metallic Materials", the entirety of which is incorporated herein by reference, and the reported results are averaged over two samples.
Charpy tests were performed in accordance with ASTM E23, "Standard Test Methods for Notched Bar Impact Testing of Metallic Materials", the entirety of which is hereby incorporated by reference, and the reported results are averaged over two or three samples. Impact tests were performed on composition 1 at temperatures of about -20°C, 0°C, and room temperature, while impact tests were performed on compositions 2 and 3 at temperatures of about -60°C, -40°C, -20°C, 0°C and room temperature.
Hardness tests were performed in accordance with ASTM E92, "Standard Test Methods for Vickers Hardness of Metallic Materials", the entirety of which is hereby incorporated by reference.

Example 1 - Continuous Cooling Transformation (CCT) behavior and microstructural evaluation of Composition 1 for cooling rates between about 0.5 and 50°C/s.

The CCT diagram derived from dilatometric measurements of composition 1 is shown in Figure 2. Illustrated in Figure 2 are traces of temperature as a function of cooling rate for transformations of about 5%, 20%, 50%, 80%, and 95%. Due to the reheating condition, about 920°C over about 10 min, the austenitic grain size prior to transformation was estimated to be about 10-20 µm, based upon the sample cooled at about 50°C/sec.
In the CCT diagram of Figure 2, two transformation regions may be observed, corresponding to cooling rates lower and higher than about 5°C/sec. For cooling rates less than about 5°/sec, the phase transformation is observed to start at about 550-600°C. The microstructure resulting under these conditions was mainly bainitic, with some retained austenite, as illustrated in the micrographs corresponding to cooling rates of about 2°C/sec and 5°C/sec in Figure 3. Under cooling rates greater than about 5°C/sec, the start transformation temperature was depressed to about 450°C, which is close to the calculated martensite transformation temperature according to the Andrews expression, about 452°C. The microstructure observed under these conditions was again mainly bainitic, as illustrated in the micrographs corresponding to cooling rates of about 10°C/sec and 20°C/sec in Figure 3. Notably, however, the bainitic structure was finer and substantially without the blocky regions of retained austenite.
Hardness measurements for composition 1 after cooling at different rates (0.5-50°C/sec) are also illustrated in CCT diagram of Figure 2. It may be observed that the hardness ranges between about 262 Hv for cooling rates of about 2°C/sec to greater than about 340 Hv for cooling rates of about 50°C.
It is further expected, based upon the high level of hardness measured at cooling rates of about 50°C/sec, that some small martensitic regions may appear at cooling rates near and above about 50°C°/sec. It should be noted, however, that substantially no large martensitic regions were observed in the microstructure corresponding to the sample cooled at about 50°C.

Example 2 - Toughness Evaluation of Composition 1 - Accelerated Cooling Condition

In order to study the impact properties of composition 1 under accelerated cooling conditions, several Charpy tests were conducted using thermal cycles discussed above with respect to the CCT diagram. Samples prepared using cooling rates of about 5°C/sec, 10°C/sec, 30°C/sec, and 45°C/sec were examined. Charpy tests were performed at temperatures of about 25°C, 0°C, and -20°C. The results of these impact tests are illustrated in Table 3 and Figures 4A and 4B, which complementarily plot impact energy (Charpy V-Notch, CVN) as a function of cooling rate and test temperature, respectively.

Table 3 - Impact Energy and Hardness of composition 1 after accelerated cooling simulations performed at Gleeble

Composition	Condition	Cooling Rate (°C/sec)	Hardness (Hv)	Impact Energy
Composition	Condition	Cooling Rate (°C/sec)	Hardness (Hv)	T(°C)	CVN (J)	Shear Area (%)
1	Accelerated Cooling	5	280	25	284	100
				0	287	91
				-20	286	89
1	Accelerated Cooling	10	305	25	335	100
				0	278	89
				-20	257	78
1	Accelerated Cooling	30	331	25	251	100
				0	229	87
				-20	240	87
1	Accelerated Cooling	45*	340	25	230	100

*Impact testing for samples cooled at about 45°C/sec was performed only for one sample at a temperature of about 25°C.

Examining Table 3 and Figures 4A and 4B, it may be observed that impact energy values fell with increasing cooling rate. Furthermore, it may be observed that excellent impact properties were obtained over the range of all cooling rates studied. For example, the impact energies measured for samples tested at temperatures between about 25°C and -20°C ranged between about 335 to 240 J. Furthermore, it can be seen from Figures 4A and 4B that the highest impact energy values correspond to samples cooled at rates ranging between about 5 to 10°C/sec. Nonetheless, even at cooling rates of about 30°C/sec, impact energy values above about 220 J were obtained at about -20°C.

Example 3 - Mechanical Evaluation of Composition 1 - As-Quenched

The tensile and impact properties of composition 1 in the as-quenched condition are illustrated in Tables 4 and 5. Table 4 - Tensile properties of composition 1 after quenching

Composition Condition YS (ksi) UTS (ksi) YS/UTS El (%)

1 As-Quenched 121 156 0.78 16

Table 5 - Impact Energy and Hardness of composition 1 after quenching

Composition Condition Hardness Charpy (10 x 10 mm)

(Hv 1 kg) T (°C) CVN (J) Shear Area (%)

1 As-Quenched 363 25 150 79

0 82 43

-20 42 24
In general, the as-quenched composition exhibited improvements in strength and impact energy over that of as-rolled samples (YS ∼ 69 ksi, UTS ∼ 99 ksi, CVN∼ 6-8 J at 25°C to -20°C). This improvement may be ascribed to a general refinement of the microstructure and the substantial disappearance of large, blocky austenitic regions.

Example 4 - Mechanical Evaluation of Composition 1 - Quenched and Tempered

a) Hardness

In order to study the tempering behavior of composition 1 in the quenched and tempered condition, samples were quenched as discussed above and tempered at temperatures ranging between about 350°C to 440°C for about 1 hr. The measured hardness values are illustrated in Figure 5. In general, it may be observed that the hardness in the as-quenched condition is about 362 Hv, falling modestly with tempering at about 300 to 400°C to within about 350 to 335 Hv. Samples tempered at about 440°C further exhibited a significant decrease in hardness, falling to about 280 +/- 20 Hv.

b) Tensile and Impact Properties

Two tempering conditions above about 400°C, 410°C and 440°C, were selected for use on quenched plates large enough to measure tensile and impact properties. Tables 6 and 7, below, summarize the experimental results along with comparable measurements for as-quenched samples.

Table 6 - Tensile properties of composition 1

Composition	Condition	Temper (°C)	Yield Strength (ksi)	UTS (ksi)	YS/UTS	El (%)
1	As-Quenched	N/A	121	156	0.78	16
1	As-Quenched Tempered	410	129	138	0.94	14
1	Quenched and Tempered	440	129	141	0.91	16

Table 7 - Impact Energy of Composition 1

Composition	Condition	Temper (°C)	Impact Energy
Composition	Condition	Temper (°C)	T(°C)	CVN (J)	Shear Area (%)
1	As-Quenched	N/A	24	150	79
			0	82	43
			-20	42	24
1	Quenched and Tempered	410	24	215	100
			0	171	74
			-20	136	59
1	Quenched and Tempered	440	24	170	84
			0	144	76
			-20	113	49

It may be observed that a good combination of strength and toughness was achieved in the quenched and tempered condition. For example, the yield and tensile strengths measured in the quenched and tempered condition were about 129 ksi and about 138-141 ksi, respectively. In contrast, the yield strength measured in the as-quenched material was less, about 121 ksi, while the tensile strength was greater, about 156 ksi.
Concurrently, the impact energies of samples in the quenched and tempered condition were found to be greater than those of samples measured at comparable temperatures in the as-quenched condition. For example, at about 24°C, samples tempered at 410 and 440°C exhibited impact energies of about 215 and 170 J, respectively, while the impact energy of the as-quenched material was about 150 J. At about -20°C, the difference in impact energies was even greater, with samples tempered at about 410 and 440°C exhibited impact energies of about 136 and 113 J, respectively, while the impact energy of the as-quenched material was about 42 J.
Without being bound by theory, it is believed that these property differences may be rationalized by the microstructure of compositions. As shown in Figure 6, in the quenched and tempered condition, the microstructure of composition 1 is bainite and martensite, with a fine dispersion of carbides, which improves the yield strength of the quenched and tempered material over that of the as-quenched material alone.

Composition 1 Summary

Examining hardness, toughness, and tensile properties in the as-quenched and quenched and tempered conditions, it may be observed that hardness decreases with increased tempering temperature, with a marked decrease beginning around 400°C. Furthermore, toughness testing in the as-quenched and the quenched and tempered conditions (410°C and 440°C) finds that toughness is generally higher in the quenched and tempered 410°C condition, as compared with the as-quenched and quenched and tempered 440°C conditions. Additionally, the yield strength exhibits modest improvement with tempering to about 410°C to 440°C while ultimate tensile strength exhibits a modest decline with tempering to about 410°C to 440°C. These results indicate that, within the range of about 410-440°C, embodiments of composition 1 in the quenched and tempered condition provide a beneficial combination of toughness and strength over the as-quenched condition alone.
Regarding accelerated cooling, excellent impact energy and hardness values were also observed. Most notably, in the range of about 10 to 20°C/s, impact energy values greater than about 220 J at about -20°C were achieved with more than about 80% ductile area. Furthermore, hardness values ranged between about 300-320 Hv.

Example 5 - Continuous Cooling Transformation (CCT) behavior and microstructural evaluation of Compositions 2-3 prior to heat treatment

The CCT diagrams derived from dilatometric measurements of compositions 2 and 3 are shown in Figures 7 and 9, respectively, for cooling rates of about 0.2, 0.5, 5, 10, 30, and 50°C/sec. The transformation start temperatures shown in these figures were determined as the first deviation from linear behavior of both dilatometric curves. The austenitic grain sizes of compositions 2 and 3 were estimated to be between about 20 to 30 µm from measurements on samples cooled at about 50°C/sec. Figures 8 and 10 further illustrate optical micrographs of compositions 2 and 3 cooled at rates of about 0.2, 0.5, 1, 10, 30, and 50°C/sec.
From the measured CCT diagrams and the observed microstructures, the transformation behavior of compositions 2 and 3 may be identified. Bainite is the main transformation product when cooling between about 5°C/sec to 30°C/sec. At lower cooling rates, polygonal ferrite is the predominant constituent. Martensite appears in composition 2 at cooling rates of about 10°C/sec and in composition 3 at about 30°C/sec and becomes the dominant phase when cooling at about 50°C/sec in both compositions.
A number of notable differences between the two compositions may also be observed. In one aspect, at cooling rates less than about 5°C/sec in composition 2, pearlite is found in large quantities, in addition to bainite. However, in composition 3, a more complex microstructure was observed, with a higher portion of bainite and some retained austenite, in addition to pearlite. Without being bound by theory, this difference may be ascribed to the lower carbon content of composition 3, which reduces the total fraction of pearlite, as well as the alloying additions of Cr and Nb, which encourage bainite formation.
In another aspect, the scale of the bainite differs between compositions 2 and 3. Despite having similar transformation temperatures and austenitic grain sizes, the bainitic structure of composition 3 is generally finer than that of composition 2. Without being bound by theory, this observation is believed to be a consequence of the Cr and Nb alloying additions.
In a further aspect, the tendency to form martensite is stronger in composition 2. As previously discussed, the lowest cooling rate at which martensite was observed in composition 2 was about 10°C/sec, while the lowest cooling rate at which martensite was observed in composition 3 was about 30°C/sec. Of further note is the observation that, while at about 30°C/sec, only a few patches of martensite appear in composition 3, the concentration of martensite in composition 2 is similar or higher to that of bainite.
From these observations, it may be understood that in low carbon content composition 3, the bainite structure is favored over a wider range of cooling rates in comparison with composition 2. Again, without being bound by theory, this observation may also be a consequence of the Cr and Nb alloying additions.
The hardness of compositions 2 and 3 values as a function of cooling rate are also illustrated in Figures 11A and 11B. Calculations performed with Creusot-Loire modeling (see Ph. Maynier, B. Jungmann, and J. Dollet, "Creusot-Loire system for the prediction of the mechanical properties of low alloy steel products", Hardenability concepts with applications to steels, Ed. D.V. Doane and J. S. Kirkaldy, The Metallurgical Society of AIME (1978), p. 518) are presented in the same graphs for comparison (dashed line). It is notable that, despite its lower carbon content, composition 3 presents a slightly higher hardness level than that of composition 2 for cooling rates below about 30°C/sec.
Without being bound by theory, this increment in hardness may be ascribed to the microstructural refinement of composition 3 already mentioned. It is also possible that, at the lower cooling rates in composition 3, some Nb dissolved during the austenization stage re-precipitates as fine carbides, which may increase hardness.

Example 6 - Mechanical Evaluation of Compositions 2-3 - As-Quenched condition

The tensile and impact properties measured for composition 3 in the as-quenched condition are presented in Tables 8 and 9 below. Hardness properties of composition 2 are also presented in Table 9. Corresponding SEM micrographs for the compositions are illustrated in Figures 12A and 12B. These results illustrate the effect of carbon on the microstructure and mechanical properties of the compositions.

Table 8 - Tensile properties of as-quenched composition 3

Composition	Condition	YS (ksi)	UTS (ksi)	YS/UTS	El (%)
3	As-Quenched	121	148	0.82	14.5

Table 9 - Impact Energy and Hardness of as-quenched compositions 2 and 3

Composition	Condition	Hardness	Charpy (10 x 10 mm)
Composition	Condition	(Hv 1Kg)	T (°C)	CVN (J)	Shear Area (%)
2	As-Quenched	350	-	-	-
3	As-Quenched	339	20	163	100
			0	161	100
			-20	162	97
			-40	96	47
			-60	67	27

The microstructure of as-quenched composition 2 was primarily martensite, with some regions of bainite (Figure 12A). Further, the hardness of composition 2 was relatively high, about 350 Hv. Based upon experience with other systems, poor toughness was expected for this system and no tensile or impact tests were performed.
In contrast, the as-quenched microstructure of composition 3 was predominantly bainitic, with small regions of martensite (Figure 12B). In this case, tensile and impact tests were carried out and a beneficial combination of properties was obtained. The yield strength was measured to be approximately 121 ksi, with a low yield strength to tensile strength ratio, about 0.82. Further, the ductile to brittle transition temperature, measured as that corresponding to about 50% shear area, was found to be about -40°C. Additionally, the impact energy was measured to be substantially constant at about 160J between about -20°C to 20°C.

Example 7 - Mechanical Evaluation of Compositions 2-3 - Quenched and tempered condition

From example 6, the as-quenched condition was found to yield beneficial properties in case of composition 3. To further probe the effect of tempering on compositions 2 and 3, similar testing and evaluation were performed on samples of compositions 2 and 3 in the quenched and tempered condition.
Figures 13A and 13B present scanning electron micrographs of the microstructure of compositions 2 and 3 in the quenched and tempered condition. In both compositions, the microstructure was composed mainly of slightly tempered bainite. There were also some small regions of tempered martensite, especially in composition 2, which possessed higher carbon content.
Hardness results from small samples of compositions 2 and 3 tempered between about 400 to 700°C are illustrated in Figure 14. It may be observed that the response of both compositions exhibit similar evolution in hardness with increasing tempering temperature. As expected, owing to its greater carbon content, composition 2 was found to exhibit greater hardness than composition 3 in the as-quenched condition and at low tempering temperatures. Conversely, however, for tempering temperatures greater than about 550°C, the hardness of composition 3 was found to be greater than that of composition 2.
Without being bound by theory, it is believed that these results may be a consequence of the Nb and Cr alloying additions to composition 3. The former may induce some precipitation hardening, while the latter may delay cementite coarsening.

In view of these observations, tensile and impact energy tests were further performed on samples of

compositions

2 and 3, which were heat treated at a temperature of about 500°C. The results of the tensile testing are given in Table 10, while impact energies are given in Table 11.

Table 10 - Tensile properties of quenched and tempered (500°C) compositions 2 and 3

Composition	Condition	YS (ksi)	UTS (ksi)	YS/UTS	El (%)
2	Quenched and Tempered	118	127	0.93	14
3	Quenched and Tempered	118	126	0.94	14

Table 11 - Impact Energy and Hardness of quenched and tempered (500°C @ 30 min) compositions 2 and 3

Composition	Condition	Hardness	Charpy (10 x 10 mm)
Composition	Condition	(Hv 1Kg)	T (°C)	CVN (J)	Shear Area (%)
2	Quenched and Tempered (500°C @ 30 min)	270	20	177	100
			0	185	100
			-20	181	100
			-40	179	100
			-60	173	97
3	Quenched and Tempered (500°C @ 30 min)	260	20	189	100
			0	189	100
			-20	189	100
			-40	175	92
			-60	143	77

The mechanical properties after tempering at about 500°C for about 30 min are found to exhibit a good combination of strength and toughness in both compositions 2 and 3. The yield strengths of compositions 2 and 3 are about 118 ksi and the ultimate tensile strengths are about 126-127 ksi.
Compositions 2 and 3 further exhibit ductile to brittle transition temperatures below about -60°C, with nearly 100% of shear area over the temperature range examined. The upper shelf energies, representing 100% of shear area, were about 180 J in both alloys, which is a good value, given the level of strength these compositions exhibit. Additionally, the compositions exhibited only minor differences in their impact energies, about 177 to 185 J in composition 2, versus about 175 to 189 J in composition 3, over the temperature range from about -40°C to 20°C.
Notably, despite their differences in alloying content, compositions 2 and 3 exhibit nearly identical tensile and impact energy properties. Without being bound by theory, in comparing the chemistries of the two compositions, it appears that the reduction in carbon content in composition 3 is approximately offset by the alloying additions of Cr and Nb.

Compositions 2 and 3 summary

Examination of compositions 2 and 3 which were air-cooled after hot rolling, then reheated and quenched, exhibited good toughness when carbon content was maintained below about 0.07% (composition 3). Furthermore, good combinations of strength and toughness were obtained when quenching and tempering at temperatures of about 500°C. In this case, excellent mechanical properties were obtained for both compositions 2 and 3, with yield strengths of about 118 ksi and impact energies of about 175-179 J at about -40°C. Additionally, almost fully ductile fracture surfaces were observed for the range of testing temperatures studied, with the ductile to brittle transition temperature well below about -60°C for both materials.

Example 7 - Simulations of thermal cycles in the heat affected zone (HAZ) of composition 2

Simulations of thermal cycles in the HAZ were performed on samples of composition 2. The Hannerz model (N.E. Hannerz, "Effect of Cb on HAZ ductility in constructional HT steels", Welding Journal, May 1975), the entirety of which is hereby incorporated by reference, was used to estimate the thermal evolution in the HAZ for different welding conditions and tube geometries. The calculated thermal cycles were reproduced at the thermo-mechanical simulator Gleeble. The microstructure and hardness of the heat treated samples were further analyzed.

These results were compared to those corresponding to a commercial, low carbon Nb-V microalloyed steel used for producing X65 heavy wall seamless tubes. The Nb-V steel possessed approximately the same carbon content and P_cm value as composition 2 but without boron addition, as illustrated in Table 12.

Table 12 - Composition of Nb-V steel and composition 2

	Nb-V	Composition 2
C (wt%)	0.09	0.09
Mn (wt%)	1.07	1.12
Si (wt%)	0.26	0.08
Ni (wt%)	0.01	0.02
Cr (wt%)	0.12	0.02
Mo (wt%)	0.10	0.26
V (wt%)	0.080	0.002
Nb (wt%)	0.021	0.003
Ti (wt%)	0.001	0.024
B (ppm)	-	13
CE(P_cm )	0.20	0.174

CE (P_{cm}) = [CE (P_{cm}) = C + Si / 30 + (Mn + Cu + Cr) / 20 + Ni / 60 + Mo / 15 + V / 10 + 5 B]

Different preheat temperatures and heat inputs were simulated for tubes of about 16 mm and/or 25 mm wall thickness (Table 13). In all cases, the maximum reheating temperature was about 1350°C in order to have the biggest possible austenitic grain size. This condition is known to be adverse for toughness because of the increased hardenability. Regarding heat inputs (HI), values between about 450 and 1210 J/mm were simulated, and preheat temperatures ranged from no preheat to approximately 250°C.

Table 13 - Welding conditions simulated at Gleeble and hardness results

#	Preheat temperature (°C)	Heat Input (J/mm)	Wall thickness (mm)	Max temp. (°C)	t85* (sec)	Cooling rate** (°C/sec)	Hardness HV1
1	150	850	16	1350	10.0	30	255±8
2	150	650	16	1350	6.0	50	281±6
3	150	450	16	1350	3.0	100	327±5
4	150	500	25	1350	3.7	81.1	294±7
5	150	500	16	1350	3.7	81.1	294±7
6	200	800	25	1350	7.5	40.0	251±8
7	200	800	16	1350	13.2	22.7	249±5
8	250	1200	25	1350	14.7	20.4	239±7
9	250	1200	16	1350	20.0	15.0	233±4
10	no preheat	550	16	1350	2.0	150.0	331±4
11	no preheat	500	16	1350	2.3	130.4	348±8
12	no preheat	500	16	1350	2.3	130.4	348±8
13	no preheat	660	16	1350	3.0	100.0	334±3
14	no preheat	940	16	1350	6.0	50.0	298±2
15	no preheat	1210	16	1350	10.0	30.0	290±7

* cooling time between 800°C and 500°C
** average cooling rate between 800°C and 500°C.

For all tests, the hardness results are presented in last column of Table 13. Assume about 300 HV as the maximum hardness, which corresponds to the maximum HAZ hardness specified by standard API 5L for X65-X80 PSL2 pipes (offshore services). From Table 13, it is clear that using a preheat of about 150°C, the minimum heat input should be about 500 J/mm for the range of analyzed wall thickness. Without the preheat, the minimum heat input would be increased to about 950 J/mm in tubes with about 25 mm wall thickness.
When comparing the hardness measured as a function of the average cooling rate with the results obtained for the Nb-V steel using the same welding conditions (Figure 15), it is clear that both steels have approximately the same hardenability. This result shows that composition 2 steel has no severe welding restrictions, because it presents substantially the same hardness behavior as a function of the cooling rate as a commercial X65 steel.
Currently there is no standard specification for the maximum hardness in the HAZ of X100 or superior grade. However, taking the maximum specified in API 5L for X80 PSL 2 (offshore) as a reference, the steel of composition 2 would meet the requirement when using about 150°C preheat and a minimum heat input of about 500 J/mm.
In summary, low carbon steels having alloying additions of boron and titanium are presented. Free nitrogen impurities are substantially consumed by reaction with titanium, forming TiN precipitates. Casting parameters are further selected so as to inhibit these precipitates from coarsening. For example, by employing cooling rates greater than about 10 to 40°C/min during casting, fine precipitates of TiN having a mean diameter less than about 50 nm may be achieved. Substantial removal of free nitrogen impurities further allows free boron to remain in solid solution, improving hardenability during austenite decomposition. These compositions may be cooled from hot rolling in air and quenched, quenching and tempering, or subjected to accelerated cooling directly after hot rolling at rates between about 5 to 50°C, yielding an excellent balance of strength and toughness.
As used throughout the specification, the term "about" should be understood to include its ordinary meaning, as understood by one of skill in the art. Where the specification uses the term "about" with respect to a particular value or range of values, the exact value or range of values provided is also contemplated as part of the disclosure.
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, changes, and/or additions 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.

Claims

A method of making a steel pipe, comprising:
providing a steel composition comprising:

0.04-0.12 wt. % carbon (C);

0.01 to 0.03 wt. % titanium (Ti);

0.0005 to 0.003 wt. % boron (B);

less than or equal to 0.008 wt. % nitrogen (N);

0.6 to 1.6 wt. % manganese (Mn);

0.05 to 0.3 wt. % silicon (Si);

less than or equal to 0.5 wt. % nickel (Ni);

less than or equal to 0.5 wt. % chromium (Cr);

less than or equal to 0.5 wt. % molybdenum (Mo);

less than or equal to 0.15 wt. % vanadium (V); and

less than or equal to 0.05 wt. % niobium (Nb);

the concentration of each element being based upon the total weight of the steel composition;

wherein 0.0005 to 0.002 wt. % boron is in solid solution for improving hardenability;

wherein substantially all of the nitrogen is present in the form TiN particles to avoid the formation of boron nitrides and achieve said boron content in solid solution, and

cooling a bar cast from the steel composition, wherein the cooling rate at about the center of the bar is selected such that TiN particles formed in the bar exhibit a mean diameter less than about 50 nm;

wherein the steel bar is cooled from casting at a rate greater than 10°C/min at about the center of the bar,

forming a pipe from the bar by reheating the steel bar to between 1200 to 1300°C;

piercing the reheated bar at temperatures ranging between 1100 to 1200°C;

and hot rolling the pierced bar at temperatures ranging between 900-1100°C; followed by cooling the steel pipe in air from hot rolling at a rate less than 1°C /sec;

and subsequently austenizing and quenching the hot rolled steel pipe; wherein hot rolling achieves an austenitic grain size of the microstructure of the steel pipe, prior to cooling from hot rolling, within a range between 20 to 50 µm, and wherein the final microstructure of the steel pipe comprises a mixture of bainite and martensite with no more than 30% of martensite if no tempering is performed after quenching and no less than 30% of martensite if tempering is performed after quenching.
The method of Claim 1, wherein the steel bar is cooled from casting at a rate greater than 30°C /min at about the center of the bar.
The method of Claim 1, wherein the elements of the steel composition are selected in concentrations such that the carbon equivalency (CE_Pcm) of the composition is less than 0.22, where CE_Pcm is calculated according to: $CE = C + \frac{Si}{30} + \frac{Mn + Cu + Cr}{30} + \frac{Ni}{60} + \frac{Mo}{15} + \frac{V}{10} + 5 B$
wherein the concentration of each element is provided in wt. %.
The method of Claim 1, further comprising tempering the quenched steel pipe at a temperature ranging between 400 to 700°C for between 10 to 60 minutes.
A method of making a steel pipe, comprising:
providing a steel composition comprising:

0.04 to 0.12 wt. % carbon (C);

0.8 to 1.6 wt. % manganese (Mn);

0.05 to 0.3 wt. % silicon (Si);

less than or equal to 0.5 wt. % nickel (Ni);

less than or equal to 0.5 wt. % chromium (Cr);

less than or equal to 0.5 wt. % molybdenum (Mo);

less than or equal to 0.15 wt. % vanadium (V);

less than or equal to 0.05 wt. % niobium (Nb);

0.01 to 0.03 wt. % titanium (Ti);

0.0005 to 0.0030 wt. % boron (B); and

less than or equal to 0.008 wt. % nitrogen (N);

wherein the concentration of each element is based upon the total weight of the steel composition and wherein 0.0005 to 0.002 wt. % boron is kept in solid solution for improving hardenability;

cooling a bar cast from the steel composition, wherein substantially all of the nitrogen in the cast steel composition is present in the form of TiN particles having a size less than 50 nm to avoid the formation of boron nitrides and achieve said boron content in solid solution;

forming a pipe from the bar by reheating the steel bar to between 1200 to 1300°C;

piercing the reheated bar at temperatures ranging between 1100 to 1200°C; and hot rolling the pierced bar at temperatures ranging between 900-1100°C; followed by

cooling the steel pipe directly after hot rolling at a rate between 5 to 50°C/sec without any subsequent heat treatment;

wherein the hot rolling achieves an austenitic grain size of the microstructure of the steel pipe, prior to cooling from hot rolling, within a range between 20 to 50 µm, and wherein the final microstructure of the steel pipe comprises a mixture of bainite and martensite with no more than 30% of martensite.
The method of Claim 5, wherein the pipe is cooled directly from hot rolling at a rate between 10 to 50°C/sec.
The method of Claim 5, wherein the pipe is cooled directly from hot rolling at a rate between 10 to 20°C/sec.
The method of Claim 5 , wherein the yield strength of the steel pipe after hot rolling and cooling, measured according to ASTM E8, is at least 100 ksi (690 MPa).
The method of Claim 5, wherein the Charpy V-notch impact energy of the pipe, strength of the steel pipe after hot rolling and cooling, measured according to ASTM E23 in full size samples is greater than 220 J at temperatures greater than or equal to -20°C.
The method of Claim 5 , wherein the composition comprises:
0.05-0.10 wt. % carbon (C);

0.8 to 1.6 wt. % manganese (Mn);

0.05 to 0.30 wt. % silicon (Si);

up to 0.4 wt. % nickel (Ni);

up to 0.3 wt. % chromium (Cr);

up to 0.3 wt. % molybdenum (Mo);

up to 0.1 wt. % vanadium (V);

up to 0.04 wt. % niobium (Nb);

0.015 to 0.025 wt. % titanium (Ti);

0.0005- 0.003 wt. % boron (B); and

less than or equal to 0.007 wt. % nitrogen (N);
A method of making a steel pipe, comprising:
providing a steel composition comprising:

0.04-0.12 wt. % carbon (C);

0.8 to 1.6 wt. % manganese (Mn);

0.05 to 0.3 wt. % silicon (Si);

less than or equal to 0.5 wt. % nickel (Ni);

less than or equal to 0.5 wt. % chromium (Cr);

less than or equal to 0.5 wt. % molybdenum (Mo);

less than or equal to 0.15 wt. % vanadium (V);

less than or equal to 0.05 wt. % niobium (Nb);

0.01 to 0.03 wt. % titanium (Ti);

0.0005-0.0030 wt. % boron (B); and

less than or equal to 0.008 wt. % nitrogen (N);

wherein the concentration of each element is based upon the total weight of the steel composition;

wherein 0.0005 to 0.002 wt. % boron is kept in solid solution for improving hardenability;

cooling a bar cast from the steel composition, wherein substantially all of the nitrogen in the cast steel composition is present in the form of TiN particles having a size less than 50 nm to avoid the formation of boron nitrides and achieve said boron content in solid solution;

forming a pipe from the bar by reheating the steel bar to between 1200 to 1300°C;

piercing the reheated bar at temperatures ranging between 1100 to 1200°C; and hot rolling the pierced bar at temperatures ranging between 900-1100°C; followed by

air cooling the steel pipe directly after hot rolling at a rate less than 1 °C/sec; and

subsequently austenizing and quenching the pipe;

wherein hot rolling achieves an austenitic grain size of the microstructure of the steel pipe, prior to cooling from hot rolling, within a range between 20 to 50 µm, and wherein the final microstructure of the steel pipe comprises a mixture of bainite and martensite with no more than 30% of martensite if no tempering is performed after quenching and no less than 30% of martensite if tempering is performed after quenching.
The method of Claim 11, wherein the steel composition comprises:
0.07 to 0.10 wt. % carbon (C);

1.0 to 1.4 wt. % manganese (Mn);

0.05 to 0.15 wt. % silicon (Si);

up to 0.4 wt. % nickel (Ni);

up to 0.35 wt. % chromium (Cr);

up to 0.3 wt. % molybdenum (Mo);

up to 0.1 wt. % vanadium (V);

up to 0.04 wt. % niobium (Nb);

0.02 to 0.03 wt. % titanium (Ti); and

0.001 to 0.002 wt. % boron (B).
The method of Claim 12, wherein the quenched steel pipe is subsequently tempered at a temperature between 400 to 600°C.
The method of Claim 13, wherein after hot rolling, cooling, austenizing and quenching, the yield strength of the pipe, measured according to ASTM E8, is greater than 100 ksi, and the Charpy V-notch impact energy of the composition, measured according to ASTM E23 in full size samples, is greater than 170 J at temperatures equal to or greater than -40°C.
The method of Claim 11 , wherein the steel composition comprises:
0.04 to 0.08 wt. % carbon;

1.0 to 1.4 wt. % manganese;

0.05 to 0.15 wt. % silicon;

up to 0.35 wt. % chromium;

0.2 to 0.3 wt. % molybdenum;

0.03 to 0.04 wt. % niobium;

0.02 to 0.03 wt. % titanium; and

0.001 to 0.002 wt. % boron.
The method of Claim 15, wherein the steel pipe is reheated in the austenitic region and quenched without subsequent tempering.
The method of Claim 16, wherein after hot rolling, cooling, austenizing and quenching, the yield strength of the pipe, measured according to ASTM E8, is greater than 100 ksi and the Charpy V-notch impact energy, measured according to ASTM E23 in full size samples, is greater than 90 J at temperatures equal to or greater than - 40°C.