WO2000039352A2 - Ultra-high strength steels with excellent cryogenic temperature toughness - Google Patents

Abstract

An ultra-high strength, weldable, low alloy steel, containing less than 9 wt% nickel, with excellent cryogenic temperature toughness in the base plate and in the heat affected zone (HAZ) when welded, having a tensile strength greater than about 830 MPa (120 ksi) and a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered fine granular bainite (FGB), or mixtures thereof, is prepared by heating a steel slab comprising iron and specified weight percentages of some or all of the additives carbon, manganese, nickel, nitrogen, copper, chromium, molybdenum, silicon, niobium, vanadium, titanium, aluminum, and boron; reducing the slab to form plate in one or more passes in a temperature range in which austenite recrystallizes; finish rolling the plate in one or more passes in a temperature range below the austenite recrystallization temperature and above the AR3 transformation temperature; quenching the finish rolled plate to a suitable Quench Stop Temperature; stopping the quenching; and tempering the plate at a suitable temperature for a period of time sufficient to cause precipitation of hardening particles.

Description

ULTRA-HIGH STRENGTH STEELS WITH EXCELLENT CRYOGENIC TEMPERATURE TOUGHNESS

FIELD OF THE INVENTION

This invention relates to ultra-high strength, weldable, low alloy steel plates with excellent cryogenic temperature toughness in both the base plate and in the heat affected zone (HAZ) when welded. Furthermore, this invention relates to a method for producing such steel plates.

BACKGROUND OF THE INVENTION Various terms are defined in the following specification. For convenience, a

Glossary of terms is provided herein, immediately preceding the claims.

Frequently, there is a need to store and transport pressurized, volatile fluids at cryogenic temperatures, i.e., at temperatures lower than about -40°C (-40°F). For example, there is a need for containers for storing and transporting pressurized liquefied natural gas (PLNG) at a pressure in the broad range of about 1035 kPa (150 psia) to about 7590 kPa (1100 psia) and at a temperature in the range of about -123°C (-190°F) to about -62°C (-80°F). There is also a need for containers for safely and economically storing and transporting other volatile fluids with high vapor pressure, such as methane, ethane, and propane, at cryogenic temperatures. For such containers to be constructed of a welded steel, the steel must have adequate strength to withstand the fluid pressure and adequate toughness to prevent initiation of a fracture, i.e., a failure event, at the operating conditions, in both the base steel and in the HAZ.

The Ductile to Brittle Transition Temperature (DBTT) delineates the two fracture regimes in structural steels. At temperatures below the DBTT, failure in the steel tends to occur by low energy cleavage (brittle) fracture, while at temperatures above the DBTT, failure in the steel tends to occur by high energy ductile fracture. Welded steels used in the construction of storage and transportation containers for the aforementioned cryogenic temperature applications and for other load-bearing, cryogenic temperature service must have DBTTs well below the service temperature in both the base steel and the HAZ to avoid failure by low energy cleavage fracture. Nickel-containing steels conventionally used for cryogenic temperature structural applications, e.g., steels with nickel contents of greater than about 3 wt%, have low DBTTs, but also have relatively low tensile strengths. Typically, commercially available 3.5 wt% Ni, 5.5 wt% Ni, and 9 wt% Ni steels have DBTTs of about -100°C (-150°F), -155°C (-250°F), and -175°C (-280°F), respectively, and tensile strengths of up to about 485 MPa (70 ksi), 620 MPa (90 ksi), and 830 MPa (120 ksi), respectively. In order to achieve these combinations of strength and toughness, these steels generally undergo costly processing, e.g., double annealing treatment. In the case of cryogenic temperature applications, industry currently uses these commercial nickel-containing steels because of their good toughness at low temperatures, but must design around their relatively low tensile strengths. The designs generally require excessive steel thicknesses for load-bearing, cryogenic temperature applications. Thus, use of these nickel-containing steels in load-bearing, cryogenic temperature applications tends to be expensive due to the high cost of the steel combined with the steel thicknesses required.

On the other hand, several commercially available, state-of-the-art, low and medium carbon high strength, low alloy (HSLA) steels, for example AISI 4320 or 4330 steels, have the potential to offer superior tensile strengths (e.g., greater than about 830 MPa (120 ksi)) and low cost, but suffer from relatively high DBTTs in general and especially in the weld heat affected zone (HAZ). Generally, with these steels there is a tendency for weldability and low temperature toughness to decrease as tensile strength increases. It is for this reason that currently commercially available, state-of-the-art HSLA steels are not generally considered for cryogenic temperature applications. The high DBTT of the HAZ in these steels is generally due to the formation of undesirable micro structures arising from the weld thermal cycles in the coarse grained and intercritically reheated HAZs, i.e., HAZs heated to a temperature of from about the Aci transformation temperature to about the Ac transformation temperature. (See Glossary for definitions of Acj and Ac₃ transformation temperatures.) DBTT increases significantly with increasing grain size and embrittling microstructural constituents, such as martensite-austenite (MA) islands, in the HAZ. For example, the DBTT for the HAZ in a state-of-the-art HSLA steel, XI 00 linepipe for oil and gas transmission, is higher than about -50°C (-60°F). There are significant incentives in the energy storage and transportation sectors for the development of new steels that combine the low temperature toughness properties of the above-mentioned commercial nickel-containing steels with the high strength and low cost attributes of the HSLA steels, while also providing excellent weldability and the desired thick section capability, i.e., the ability to provide substantially the desired microstructure and properties (e.g., strength and toughness), particularly in thicknesses equal to or greater than about 25 mm (1 inch).

In non-cryogenic applications, most commercially available, state-of-the-art, low and medium carbon HSLA steels, due to their relatively low toughness at high strengths, are either designed at a fraction of their strengths or, alternatively, processed to lower strengths for attaining acceptable toughness. In engineering applications, these approaches lead to increased section thickness and therefore, higher component weights and ultimately higher costs than if the high strength potential of the HSLA steels could be fully utilized. In some critical applications, such as high performance gears, steels containing greater than about 3 wt% Ni (such as AISI 48XX, SAE 93XX, etc.) are used to maintain sufficient toughness. This approach leads to substantial cost penalties to access the superior strength of the HSLA steels. An additional problem encountered with use of standard commercial HSLA steels is hydrogen cracking in the HAZ, particularly when low heat input welding is used.

There are significant economic incentives and a definite engineering need for low cost enhancement of toughness at high and ultra-high strengths in low alloy steels. Particularly, there is a need for a reasonably priced steel that has ultra-high strength, e.g., tensile strength greater than about 830 MPa (120 ksi), and excellent cryogenic temperature toughness, e.g. DBTT lower than about -62°C (-80°F), both in the base plate when tested in the transverse direction (see Glossary for definition of transverse direction) and in the HAZ, for use in commercial cryogenic temperature applications. Consequently, the primary objects of the present invention are to improve the state-of-the-art high strength, low alloy steel technology for applicability at cryogenic temperatures in three key areas: (i) lowering of the DBTT to less than about -62°C (-80°F) in the base steel in the transverse direction and in the weld HAZ, (ii) achieving tensile strength greater than 830 MPa (120 ksi), and (iii) providing superior weldability. Other objects of the present invention are to achieve the aforementioned HSLA steels with thick section capability, preferably, for thicknesses equal to or greater than about 25 mm (1 inch) and to do so using current commercially available processing techniques so that use of these steels in commercial cryogenic temperature processes is economically feasible.

SUMMARY OF THE INVENTION

Consistent with the above-stated objects of the present invention, a processing methodology is provided wherein a low alloy steel slab of the desired chemistry is reheated to an appropriate temperature, hot rolled to form steel plate, rapidly cooled, at the end of hot rolling, by quenching with a suitable fluid, such as water, to a suitable Quench Stop Temperature (QST), to transform the microstructure of the steel preferably to predominantly fine-grained lath martensite, fine-grained lower bainite, fine granular bainite (FGB) or mixtures thereof, or, more preferably, to substantially 100% fine-grained lath martensite, and then tempered within a suitable temperature range to produce a microstructure in the tempered steel preferably comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered FGB, or mixtures thereof, or, more preferably comprising substantially 100% tempered fine-grained lath martensite. The FGB of the present invention is an aggregate comprising bainitic ferrite as a major constituent (at least about 50 vol%) and particles of mixtures of martensite and retained austenite as minor constituents (less than about 50 vol%). As used in describing the present invention, quenching refers to accelerated cooling by any means whereby a fluid selected for its tendency to increase the cooling rate of the steel is utilized, as opposed to air cooling the steel to ambient temperature. In one embodiment of this invention, the steel plate is air cooled to ambient temperature after quenching is stopped and prior to tempering. As used in describing the present invention, and in the claims, "predominantly", "predominant" and "major" all mean at least about 50 volume percent and "minor" means less than about 50 vol%.

A steel slab processed according to this invention is manufactured in a customary fashion and, in one embodiment, comprises iron and the following alloying elements, preferably in the weight ranges indicated in the following Table I:

Table I

Alloying Element Range (wt%)

carbon (C) 0.03 - 0.12, more preferably 0.03 - 0.07 manganese (Mn) up to 2.5, more preferably 1.0 - 1.8 nickel (Ni) 1.0 - 3.0, more preferably 1.5 - 3.0 copper (Cu) up to 1.5, more preferably up to 1.0 molybdenum (Mo) up to 0.8, more preferably 0J - 0.5 niobium (Nb) 0.02 - 0.1, more preferably 0.03 - 0.1 titanium (Ti) 0.008 - 0.03, more preferably 0.01 - 0.02 aluminum (Al) 0.001 - 0.05, more preferably 0.005 - 0.03 nitrogen (N) 0.002 - 0.005, more preferably 0.002 - 0.003

Vanadium (V) is sometimes added to the steel, preferably up to about 0J0 wt%, and more preferably about 0.02 wt% to about 0J wt%.

Chromium (Cr) is sometimes added to the steel, preferably up to about 1.0 wt%, and more preferably about 0J wt% to about 0.6 wt%.

Silicon (Si) is sometimes added to the steel, preferably up to about 0.5 wt%, more preferably about 0.01 wt% to about 0.5 wt%, and even more preferably about 0.05 wt% to about 0.1 wt%.

Boron (B) is sometimes added to the steel, preferably up to about 0.0020 wt%, and more preferably about 0.0006 wt% to about 0.0015 wt%.

The steel preferably contains at least about 1 wt% nickel. Nickel content of the steel can be increased above about 3 wt% if desired to enhance performance after welding. Each 1 wt% addition of nickel is expected to lower the DBTT of the steel by about 10°C (18°F). Nickel content is preferably less than 9 wt%, more preferably less than about 6 wt%. Nickel content is preferably minimized in order to minimize cost of the steel. If nickel content is increased above about 3 wt%, manganese content can be decreased below about 0.5 wt% down to 0.0 wt%.

Additionally, residuals are preferably substantially minimized in the steel. Phosphorous (P) content is preferably less than about 0.01 wt%. Sulfur (S) content is preferably less than about 0.004 wt%. Oxygen (O) content is preferably less than about 0.002 wt%. The specific microstructure obtained in this invention is dependent upon both the chemical composition of the alloy steel slab that is processed and the actual processing steps that are followed in processing the steel. For example, without hereby limiting this invention, some specific microstructures that are obtained are as follows. In one embodiment, a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, or mixtures thereof is produced. The other constituents in this embodiment may comprise tempered versions of FGB, polygonal ferrite (PF), deformed ferrite (DF), acicular ferrite (AF), upper bainite (UB), degenerate upper bainite (DUB) and the like, all as are familiar to those that are skilled in the art. Also, depending upon the tempering temperature, up to about 10 vol% of retained austenite or reverted austenite, as are familiar to those skilled in the art, may be present in the microstructure. This embodiment provides tensile strengths generally exceeding about 930 MPa (135 ksi). In yet another embodiment of this invention, the steel plate has a microstructure comprising predominantly tempered FGB. The other constituents that comprise the microstructure may include tempered versions of fine-grained lath martensite, fine-grained lower bainite, PF, DF, AF, UB, DUB, retained austenite, reverted austenite, or the like. This embodiment provides tensile strengths generally in the lower range of this invention, i.e., tensile strengths of about and greater than 830 MPa (120 ksi). As is discussed in greater detail herein, the value of NC, a factor defined by the chemistry of the steel (as further discussed herein and in the Glossary), also impacts the strength and thick section capability, as well as microstructure, of steels according to this invention. Also, consistent with the above- stated objects of the present invention, steels processed according to the present invention are especially suitable for many cryogenic temperature applications in that the steels have the following characteristics, preferably, without thereby limiting the invention, for steel plate thicknesses of about 25 mm (1 inch) and greater: (i) DBTT lower than about -62°C (-80°F), preferably lower than about -73°C (-100°F), more preferably lower than about -100°C (-150°F) and even more preferably lower than about -123°C (-190°F) in the base steel in the transverse direction and in the weld HAZ, (ii) tensile strength greater than about 830 MPa (120 ksi), preferably greater than about 860 MPa (125 ksi), more preferably greater than about 900 MPa (130 ksi) and even more preferably greater than about 1000 MPa (140 ksi), (iii) superior weldability, and (iv) improved toughness over standard, commercially available, HSLA steels.

DESCRIPTION OF THE DRAWINGS The advantages of the present invention will be better understood by referring to the following detailed description and the attached drawings in which:

FIG. 1A is a schematic illustration of austenite grain size in a steel slab after reheating according to the present invention;

FIG. IB is a schematic illustration of prior austenite grain size (see Glossary) in a steel slab after hot rolling in the temperature range in which austenite recrystallizes, but prior to hot rolling in the temperature range in which austenite does not recrystallize, according to the present invention; and

FIG. 1C is a schematic illustration of the elongated, pancake grain structure in austenite, with very fine effective grain size in the through-thickness direction, of a steel plate upon completion of rolling in TMCP according to the present invention.

While the present invention will be described in connection with its preferred embodiments, it will be understood that the invention is not limited thereto. On the contrary, the invention is intended to cover all alternatives, modifications, and equivalents which may be included within the spirit and scope of the invention, as defined by the appended claims. DETAILED DESCRIPTION OF THE INVENTION

The present invention relates to the development of new HSLA steels meeting the above-described challenges. The invention is based on a novel combination of steel chemistry and processing for providing both intrinsic and microstructural toughening to lower DBTT as well as to enhance toughness at high tensile strengths. Intrinsic toughening is achieved by the judicious balance of critical alloying elements in the steel as described in detail in this specification. Microstructural toughening results from achieving a very fine effective grain size as well as producing fine-grained martensitic and lower bainitic laths occurring in fine packets and/or FGB with a mean grain dimension much finer than the prior austenite pancakes.

Additionally, in the present invention, dispersion strengthening from fine copper-rich precipitates and/or mixed carbides/ carbonitrides is utilized to optimize strength and toughness during the tempering of the lath martensitic/lower bainitic/FGB structure. Fine effective grain size is accomplished in two ways in the present invention. First, thermo-mechanical controlled rolling processing ("TMCP") as described in detail hereinafter is used to establish fine pancake structure in austenite at the end of rolling in the TMCP processing. This is an important first step in the overall refinement of microstructure in the present invention. Second, further refinement of pancakes is achieved through transformation of austenite pancakes to packets of lath martensite/lower bainite, grains of FGB or mixtures thereof. Fine-grained lath martensite and fine-grained lower bainite occur in packets within the austenite pancakes consisting of several similarly oriented laths. Typically, there is more than one packet in a pancake and the packet itself is made up of about 5 to 8 laths. Adjacent packets are separated by high angle boundaries. The packet width is the effective grain size in these structures and it has a significant effect on the cleavage fracture resistance and the DBTT, with finer packet widths providing lower DBTT. In the present invention, the preferred mean packet width is less than about 5 microns, and more preferably, less than about 3 microns and even more preferably less than about 2 microns. (See Glossary for definition of "high angle boundary") The FGB of the present invention is an aggregate comprising bainitic ferrite as a major constituent and particles of mixtures of martensite and retained austenite as minor constituents. The FGB of the present invention has a very fine grain size mimicking the mean packet width of the fine-grained lath martensite and fine-grained lower bainite microstructures described above. The FGB can form during the quenching to the QST and/or during the slow cooling to ambient from the QST in the steels of the present invention especially at the center of a thick, > 25 mm, plate when the total alloying in the steel is low and/or if the steel does not have sufficient "effective" boron, that is, boron that is not tied up in oxide and/or nitride. In these instances and depending on the cooling rate for the quenching and the overall plate chemistry, FGB may form either as a minor or as a predominant constituent. In the present invention, the preferred mean grain size of the FGB is less than about 3 microns, more preferably less than about 2 microns, even more preferably less than about 1 micron. Adjacent grains of the bainitic ferrite form high angle boundaries in which the grain boundary separates two adjacent grains whose crystallographic orientations differ typically by more than 15°, whereby these boundaries are quite effective in crack deflection and in enhancing crack tortuosity. In the FGB of the present invention the martensite is preferably of a low carbon (< 0.4 wt%), dislocated type with little or no twinning. The vol% of these minor constituents in the FGB of the present invention can vary depending on the steel composition and processing but is preferably less than about 40 vol%, more preferably less than about 20 vol%, and even more preferably less than about 10 vol% of the FGB. The martensite/retained austenite particles of the FGB are effective in providing additional crack deflection and tortuosity within the FGB. It has been found in this invention that, for carbon contents in the steel of about 0.030 to 0.065 wt%, the amount of FGB (averaged over thickness) in the microstructure is preferably limited to less than about 40 vol% in order for the strength of the plate exceed about 930 MPa (135 ksi). As used in describing this invention, "effective grain size" refers to mean austenite pancake thickness upon completion of rolling in the TMCP processing according to this invention and to mean packet width or grain size upon completion of transformation of the austenite pancakes to packets of lath martensite/lower bainite or FGB, respectively. As is further discussed hereinafter, D'" in FIG. 1C, illustrates pancake thickness upon completion of rolling in TMCP processing according to this invention. Packets and grains of FGB form inside of the pancakes. Packet width or FGB grain size are not illustrated in the drawings. This integrated approach provides for a very fine effective grain size, especially in the through thickness direction of a steel plate according to this invention.

In accordance with the foregoing, a method is provided for preparing an ultra- high strength steel plate having a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered FGB or mixtures thereof, wherein the method comprises the steps of (a) heating a steel slab to a reheating temperature sufficiently high to (i) substantially homogenize the steel slab, (ii) dissolve substantially all carbides and carbonitrides of niobium and vanadium in the steel slab, and (iii) establish fine initial austenite grains in the steel slab; (b) reducing the steel slab to form steel plate in one or more hot rolling passes in a first temperature range in which austenite recrystallizes; (c) further reducing the steel plate in one or more hot rolling passes in a second temperature range below about the T^ temperature and above about the Ar₃ transformation temperature; (d) quenching the steel plate at a cooling rate of at least about 10°C per second (18°F/sec) to a Quench Stop Temperature below about 600°C (1112°F), more preferably below about the M_s transformation temperature plus 200°C (360°F); (e) stopping the quenching; and (f) tempering the steel plate at a tempering temperature from about 400°C (752°F) up to about the Ac^*ι transformation temperature, preferably up to, but not including, the Aci transformation temperature, for a period of time sufficient to cause precipitation of hardening particles, i.e., one or more of ε-copper, Mo₂C, or the carbides and carbonitrides of niobium and vanadium. In some embodiments, cementite (iron-rich carbide) may also be precipitated upon tempering. The period of time sufficient to cause precipitation of hardening particles depends primarily on the thickness of the steel plate, the chemistry of the steel plate, and the tempering temperature, and can be determined by one skilled in the art. (See Glossary for definitions of predominantly, of hardening particles, of T_nr temperature, of Ar₃, M_s, and Aci transformation temperatures, and of Mo₂C). The QST can be ambient temperature or below. In one embodiment, the method of this invention further comprises the step of allowing the steel plate to air cool to ambient temperature from the QST before the step of tempering. In another embodiment, the steel plate can be directly tempered following the completion of quenching to the QST. To ensure high strength of greater than about 930 MPa (135 ksi) and ambient and cryogenic temperature toughness, steels according to this invention preferably have a microstructure comprised of predominantly tempered fine-grained lower bainite, tempered fine-grained lath martensite, or mixtures thereof. As used in describing the present invention, and in the claims, "predominantly" means at least about 50 volume percent. More preferably, the microstructure comprises at least about 60 volume percent to about 80 volume percent tempered fine-grained lower bainite, tempered fine-grained lath martensite, or mixtures thereof. Even more preferably, the microstructure comprises at least about 90 volume percent tempered fine-grained lower bainite, tempered fine-grained lath martensite, or mixtures thereof. Most preferably, the microstructure comprises substantially 100% tempered fine-grained lath martensite. The remainder of the microstructure may comprise tempered versions of FGB, PF, DF, AF, UB, DUB, retained austenite, reverted austenite, or the like. For lower strengths, i.e., less than about 930 MPa (135 ksi) but higher than about 830 MPa (120 ksi), the steel may have a microstructure comprising predominantly tempered FGB. The remainder of the microstructure may comprise tempered versions of fine-grained lath martensite, fine-grained lower bainite, PF, DF, AF, UB, DUB, retained austenite, reverted austenite, or the like. In all the embodiments of the present invention, it is preferable to substantially minimize (to less than about 10 vol%, more preferably less than about 5 vol% of the microstructure) the formation of embrittling constituents such as upper bainite, twinned martensite and MA.

One embodiment of this invention includes a method for preparing a steel plate having a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, or mixtures thereof, said method comprising the steps of: (a) heating a steel slab to a reheating temperature sufficiently high to (i) substantially homogenize said steel slab, (ii) dissolve substantially all carbides and carbonitrides of niobium and vanadium in said steel slab, and (iii) establish fine initial austenite grains in said steel slab; (b) reducing said steel slab to form steel plate in one or more hot rolling passes in a first temperature range in which austenite recrystallizes; (c) further reducing said steel plate in one or more hot rolling passes in a second temperature range below about the T_m temperature and above about the Ar₃ transformation temperature; (d) quenching said steel plate at a cooling rate of about 10°C per second to about 40°C per second (18°F/sec - 72°F/sec) to a

Quench Stop Temperature below about the M_s transformation temperature plus 200°C

(360°F); (e) stopping said quenching; and (f) tempering said steel plate at a tempering temperature from about 400°C (752°F) to about the Aci transformation temperature for a period of time sufficient to cause precipitation of hardening particles, said steps being performed so as to facilitate transformation of said microstructure of said steel plate to predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, or mixtures thereof.

Processing of the Steel Slab

(1) Lowering of DBTT

Achieving a low DBTT, e.g., lower than about -62°C (-80°F), in the transverse direction of the base plate and in the HAZ is a key challenge in the development of new HSLA steels for cryogenic temperature applications. The technical challenge is to maintain increase the strength in the present HSLA technology while lowering the DBTT, especially in the HAZ. The present invention utilizes a combination of alloying and processing to alter both the intrinsic as well as microstructural contributions to fracture resistance in a way to produce a low alloy steel with excellent cryogenic temperature properties in the base plate and in the HAZ, as hereinafter described.

In this invention, microstructural toughening is exploited for lowering the base steel DBTT. A key component of this microstructural toughening consists of refining prior austenite grain size and modifying the grain morphology, aimed at enhancing the interfacial area of the high angle boundaries per unit volume in the steel plate. As is familiar to those skilled in the art, "grain" as used herein means an individual crystal in a polycrystalline material, and "grain boundary" as used herein means a narrow zone in a metal corresponding to the transition from one crystallographic orientation to another, thus separating one grain from another. As used herein, a "high angle grain boundary" is a grain boundary that separates two adjacent grains whose crystallographic orientations differ by more than about 8°. Also, as used herein, a "high angle boundary" is a boundary that effectively behaves as a high angle grain boundary, i.e., a boundary that tends to deflect a propagating crack or fracture and, thus, induces tortuosity in a fracture path.

The contribution from fhermo-mechanical controlled rolling processing (TMCP) to the total interfacial area of the high angle boundaries per unit volume, Sv , is defined by the following equation:

Sv = - i f 1 + R + - ιλ + 0.63(r - 30) d V RJ

where: d is the average austenite grain size in a hot-rolled steel plate prior to rolling in the temperature range in which austenite does not recrystallize (prior austenite grain size);

R is the reduction ratio (original steel slab thickness/final steel plate thickness); and r is the percent reduction in thickness of the steel due to hot rolling in the temperature range in which austenite does not recrystallize.

It is well known in the art that as the Sv of a steel increases, the DBTT decreases, due to crack deflection and the attendant tortuosity in the fracture path at the high angle boundaries. In commercial TMCP practice, the value of R is fixed for a given plate thickness and the upper limit for the value of r is typically 75. Given fixed values for R and r , Sv can only be substantially increased by decreasing d , as evident from the above equation. To decrease d in steels according to the present invention, Ti-Nb microalloying is used in combination with optimized TMCP practice. For the same total amount of reduction during hot rolling/deformation, a steel with an initially finer average austenite grain size will result in a finer finished average austenite grain size. Therefore, in this invention the amount of Ti-Nb additions are optimized for low reheating practice while producing the desired austenite grain growth inhibition during TMCP. Referring to FIG. 1 A, a relatively low reheating temperature, preferably between about 955°C and about 1100°C (1750°F - 2012°F), is used to obtain initially an average austenite grain size D' of less than about 120 microns in reheated steel slab 10' before hot deformation. Processing according to this invention avoids the excessive austenite grain growth that results from the use of higher reheating temperatures, i.e., greater than about 1100°C (2012°F), in conventional TMCP. To promote dynamic recrystallization induced grain refining, heavy per pass reductions greater than about 10% are employed during hot rolling in the temperature range in which austenite recrystallizes. Referring now to FIG. IB, processing according to this invention provides an average prior austenite grain size D" (i.e., d ) of less than about 50 microns, preferably less than about 30 microns, more preferably less than about 20 microns, and even more preferably less than about 10 microns, in steel slab 10" after hot rolling (deformation) in the temperature range in which austenite recrystallizes, but prior to hot rolling in the temperature range in which austenite does not recrystallize. Additionally, to produce an effective grain size reduction in the through-thickness direction, heavy reductions, preferably exceeding about 70% cumulative, are carried out in the temperature range below about the T_m temperature but above about the Ar₃ transformation temperature.

Referring now to FIG. 1C, TMCP according to this invention leads to the formation of an elongated, pancake structure in austenite in a finish rolled steel plate 10'" with very fine effective grain size D"' in the through-thickness direction, e.g., effective grain size D'" less than about 10 microns, preferably less than about 8 microns, and even more preferably less than about 5 microns, and yet more preferably less than about 3 microns, thus enhancing the interfacial area of high angle boundaries, e.g., 11, per unit volume in steel plate 10'", as will be understood by those skilled in the art. To minimize anisotropy in mechanical properties in general and to enhance the toughness and DBTT in the transverse direction, it is helpful to minimize the austenite pancake aspect ratio, that is, the mean ratio of pancake length to pancake thickness. In the present invention through the control of the TMCP processing parameters as described above, the aspect ratio for the pancakes is kept preferably less than about 100, more preferably less than about 75, even more preferably less than about 50 and yet more preferably less than about 25. In somewhat greater detail, a steel according to this invention is prepared by forming a slab of the desired composition as described herein; heating the slab to a temperature of from about 955°C to about 1100°C (1750°F - 2012°F), preferably from about 955°C to about 1065°C (1750°F - 1950°F); hot rolling the slab to form steel plate in one or more passes providing about 30 percent to about 70 percent reduction in a first temperature range in which austenite recrystallizes, i.e., above about the T_m temperature, and further hot rolling the steel plate in one or more passes providing about 40 percent to about 80 percent reduction in a second temperature range below about the T^- temperature and above about the Ar transformation temperature. The hot rolled steel plate is then quenched at a cooling rate of at least about 10°C per second (18°F/sec ) to a suitable QST below about 600°C (1112°F), preferably below about the M_s transformation temperature plus 200°C (360°F), at which time the quenching is terminated. The cooling rate for the quenching to the QST is preferably faster than about 10°C per second (18°F/sec) and even more preferably faster than about 20°C per second (36°F/sec). Without hereby limiting this invention, the cooling rate in one embodiment of this invention is about 10°C per second to about 40°C per second (18°F/sec - 72°F/sec). In one embodiment of this invention, the steel plate is then air cooled to ambient temperature. This processing is used to produce a microstructure preferably comprising predominantly fine-grained lath martensite, fine-grained lower bainite, FGB or mixtures thereof, or, more preferably comprising substantially 100% fine-grained lath martensite.

The thus direct quenched martensite, lower bainite and/or FGB in steels according to this invention have high strength but their toughness can be improved by tempering at a suitable temperature from above about 400°C (752°F) up to about the Aci transformation temperature. Tempering of steel within this temperature range also leads to reduction of the quenching stresses which in turn leads to enhanced toughness. While tempering can enhance the toughness of the steel, it normally leads to substantial loss of strength. In the present invention, the usual strength loss from tempering is offset by inducing precipitate dispersion hardening. Dispersion hardening from fine copper precipitates and/or mixed carbides / carbonitrides are utilized to optimize strength and toughness during the tempering of the martensitic/bainitic/FGB structure. The unique chemistry of the steels of this invention allows for tempering within the broad range of about 400°C to about 650°C (750°F - 1200°F) without any significant loss of the as-quenched strength. The steel plate is preferably tempered at a tempering temperature from above about 400°C (752°F) to below the Ac^*ι transformation temperature for a period of time sufficient to cause precipitation of hardening particles (as defined herein). This processing facilitates transformation of the microstructure of the steel plate to predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered FGB or mixtures thereof. Again, the period of time sufficient to cause precipitation of hardening particles depends primarily on the thickness of the steel plate, the chemistry of the steel plate, and the tempering temperature, and can be determined by one skilled in the art.

As is understood by those skilled in the art, as used herein percent reduction in thickness refers to percent reduction in the thickness of the steel slab or plate prior to the reduction referenced. For purposes of explanation only, without thereby limiting this invention, a steel slab of about 254 mm (10 inches) thickness may be reduced about 50% (a 50 percent reduction), in a first temperature range, to a thickness of about 127 mm (5 inches) then reduced about 80% (an 80 percent reduction), in a second temperature range, to a thickness of about 25 mm (1 inch). As used herein, "slab" means a piece of steel having any dimensions.

The steel slab is preferably heated by a suitable means for raising the temperature of substantially the entire slab, preferably the entire slab, to the desired reheating temperature, e.g., by placing the slab in a furnace for a period of time. The specific reheating temperature that should be used for any steel composition within the range of the present invention may be readily determined by a person skilled in the art, either by experiment or by calculation using suitable models. Additionally, the furnace temperature and reheating time necessary to raise the temperature of substantially the entire slab, preferably the entire slab, to the desired reheating temperature may be readily determined by a person skilled in the art by reference to standard industry publications. Except for the reheating temperature, which applies to substantially the entire slab, subsequent temperatures referenced in describing the processing method of this invention are temperatures measured at the surface of the steel. The surface temperature of steel can be measured by use of an optical pyrometer, for example, or by any other device suitable for measuring the surface temperature of steel. The cooling rates referred to herein are those at the center, or substantially at the center, of the plate thickness; and the Quench Stop Temperature (QST) is the highest, or substantially the highest, temperature reached at the surface of the plate, after quenching is stopped, because of heat transmitted from the mid-thickness of the plate. For example, during processing of experimental heats of a steel composition according to this invention, a thermocouple is placed at the center, or substantially at the center, of the steel plate thickness for center temperature measurement, while the surface temperature is measured by use of an optical pyrometer. A correlation between center temperature and surface temperature is developed for use during subsequent processing of the same, or substantially the same, steel composition, such that center temperature may be determined via direct measurement of surface temperature. Also, the required temperature and flow rate of the quenching fluid to accomplish the desired accelerated cooling rate may be determined by one skilled in the art by reference to standard industry publications.

For any steel composition within the range of the present invention, the temperature that defines the boundary between the recrystallization range and non-recrystallization range, the T_m temperature, depends on the chemistry of the steel, particularly the carbon concentration and the niobium concentration, on the reheating temperature before rolling, and on the amount of reduction given in the rolling passes. Persons skilled in the art may determine this temperature for a particular steel according to this invention either by experiment or by model calculation. Similarly, the Aci, Ar , and M_s transformation temperatures referenced herein may be determined by persons skilled in the art for any steel according to this invention either by experiment or by model calculation.

Although the microstructural approaches described above are useful for lowering DBTT in the base steel plate, they are not fully effective for maintaining sufficiently low DBTT in the coarse grained regions of the weld HAZ. Thus, the present invention provides a method for maintaining sufficiently low DBTT in the coarse grained regions of the weld HAZ by utilizing intrinsic effects of alloying elements, as described in the following.

Leading ferritic cryogenic temperature steels are generally based on body-centered cubic (BCC) crystal lattice. While this crystal system offers the potential for providing high strengths at low cost, it suffers from a steep transition from ductile to brittle fracture behavior as the temperature is lowered. This can be fundamentally attributed to the strong sensitivity of the critical resolved shear stress (CRSS) (defined herein) to temperature in BCC systems, wherein CRSS rises steeply with a decrease in temperature thereby making the shear processes and consequently ductile fracture more difficult. On the other hand, the critical stress for brittle fracture processes such as cleavage is less sensitive to temperature. Therefore, as the temperature is lowered, cleavage becomes the favored fracture mode, leading to the onset of low energy brittle fracture. The CRSS is an intrinsic property of the steel and is sensitive to the ease with which dislocations can cross slip upon deformation; that is, a steel in which cross slip is easier will also have a low CRSS and hence a low

DBTT. Some face-centered cubic (FCC) stabilizers such as Ni are known to promote cross slip, whereas BCC stabilizing alloying elements such as Si, Al, Mo, Nb and V discourage cross slip. In the present invention, content of FCC stabilizing alloying elements, such as Ni and Cu, is preferably optimized, taking into account cost considerations and the beneficial effect for lowering DBTT, with Ni alloying of preferably at least about 1.0 wt% and more preferably at least about 1.5 wt%; and the content of BCC stabilizing alloying elements in the steel is substantially minimized.

As a result of the intrinsic and microstructural toughening that results from the unique combination of chemistry and processing for steels according to this invention, the steels have excellent cryogenic temperature toughness in both the base plate and the HAZ after welding. DBTTs in both the base plate in the transverse direction and the HAZ after welding of these steels are lower than about -62°C (-80°F) and can be lower than about -107°C (-160°F). (2) Tensile Strength greater than 830 MPa (120 ksi) and Thick Section Capability

Generally, upon tempering, plain carbon and low alloy martensitic/bainitic and FGB steels with no strong carbide formers soften or lose their as-quenched strength, the degree of this strength loss being a function of the specific chemistry of the steel and of the tempering temperature and duration. In the steels of the present invention, the loss in strength during tempering is substantially ameliorated by fine precipitation of hardening particles. The unique chemistry of the steels of this invention allows for tempering within the broad range of about 400°C to about 650°C (750°F - 1200°F) without any significant loss of the as-quenched strength. Within this broad tempering range, strengthening results from hardening particle precipitation occurring or peaking at various temperature regimes; i.e., within this broad range, sufficient precipitation of hardening particles occurs to provide cumulative strength adequate to compensate for the loss of strength normally associated with tempering. The processing flexibility provided by the ability to temper within this broad range is advantageous.

In the present invention, the desired strength is obtained at a relatively low carbon content with the attendant advantages in weldability and excellent toughness in both the base steel and in the HAZ. A minimum of about 0.03 wt% C is preferred in the overall alloy for attaining tensile strength greater than 830 MPa (120 ksi).

While alloying elements, other than C, Cu, Cr, Mo, V, Nb, and Ti, in tempered steels according to this invention are substantially inconsequential as regards the maximum attainable strength in the tempered steel, these elements are desirable to provide the required thick section capability and strength for plate thickness equal to or greater than about 25 mm (1 inch) and for a range of cooling rates desired for processing flexibility. This is important as the actual cooling rate at the mid section of a thick plate is lower than that at the surface. The microstructure of the surface and center can thus be quite different unless the steel is designed to eliminate its sensitivity to the difference in cooling rate between the surface and the center of the plate. In this regard, Mn and Mo alloying additions, and especially the combined additions of Mn, Mo and B, are particularly effective. In the present invention, these additions are optimized for hardenability, weldability, low DBTT and cost considerations. As stated previously in this specification, from the point of view of lowering DBTT, it is essential that the total BCC alloying additions be kept to a minimum. The preferred chemistry targets and ranges are set to meet these and the other requirements of this invention. In order to achieve the strength and thick section capability of the steels of this invention for plate thicknesses equal to or greater than about 25 mm, the Nc, a factor defined by the chemistry of the steel as shown below, is preferably in the range of about 2.5 to about 4.0 for steels with effective B additions, and is preferably in the range of about 3.0 to about 4.5 for steels with no added B. More preferably, for B containing steels according to this invention Nc is preferably greater than about 2.8, even more preferably greater than about 3.0. For steels according to this invention without added B, Nc preferably is greater than about 3J and even more preferably greater than about 3.5. Generally steels with Nc in the high end of the preferred range, that is, greater than about 3.0 for steels with effective B additions and 3.5 for steels without added B, of this invention when processed according to the objects of this invention result in a microstructure comprising, predominantly, tempered fine-grained lower bainite, tempered fine-grained lath martensite, or mixtures thereof. On the other hand, steels with Nc in the lower end of the preferred range shown above tend to form a predominantly tempered FGB microstructure.

Nc = 12.0*C + Mn + 0.8*Cr + 0.15*(Ni + Cu)+ 0.4*Si + 2.0*V + 0.7* Nb + 1.5*Mo, where C, Mn, Cr, Ni, Cu, Si, V, Nb, Mo represent their respective wt% in the steel.

(3) Superior Weldability For Low Heat Input Welding

The steels of this invention are designed for superior weldability. The most important concern, especially with low heat input welding, is cold cracking or hydrogen cracking in the coarse grained HAZ. It has been found that for steels of the present invention, cold cracking susceptibility is critically affected by the carbon content and the type of HAZ microstructure, not by the hardness and carbon equivalent, which have been considered to be the critical parameters in the art. In order to avoid cold cracking when the steel is to be welded under no or low preheat (lower than about 100°C (212°F)) welding conditions, the preferred upper limit for carbon addition is about 0J wt%. As used herein, without limiting this invention in any aspect, "low heat input welding" means welding with arc energies of up to about 2.5 kilojoules per millimeter (kJ/mm) (7.6 kJ/inch). Lower bainite or auto-tempered lath martensite microstructures offer superior resistance to cold cracking. Other alloying elements in the steels of this invention are carefully balanced, commensurate with the hardenability and strength requirements, to ensure the formation of these desirable microstructures in the coarse grained HAZ.

Role of Alloying Elements in the Steel Slab

The role of the various alloying elements and the preferred limits on their concentrations for the present invention are given below:

Carbon (C) is one of the most effective strengthening elements in steel. It also combines with the strong carbide formers in the steel such as Ti, Nb, V and Mo to provide grain growth inhibition and precipitation strengthening during tempering. Carbon also enhances hardenability, i.e., the ability to form harder and stronger microstructures in the steel during cooling. If the carbon content is less than about 0.03 wt%, it is not sufficient to induce the desired strengthening, viz., greater than about 830 MPa (120 ksi) tensile strength, in the steel. If the carbon content is greater than about 0J2 wt%, the steel will be susceptible to cold cracking during welding and the toughness is reduced in the steel plate and its HAZ on welding. Carbon content in the range of about 0.03 wt% to about 0J2 wt% is preferred to produce the desired strength and HAZ microstructures, viz., auto-tempered lath martensite and lower bainite. Even more preferably, the upper limit for carbon content is about 0.07 wt%. Manganese (Mn) is a matrix strengthener in steels and also contributes strongly to the hardenability. Mn is a key, inexpensive alloying addition to promote fine-grained lath martensite/fme-grained lower bainite and to prevent excessive FGB in thick section, > 25 mm, plates which can lead to reduction in strength. A minimum amount of 0.5 wt% Mn is preferred for achieving the desired high strength in plate thickness exceeding about 25 mm (1 inch), and a minimum of at least about 1.0 wt% Mn is even more preferred. However, too much Mn can be harmful to toughness, so an upper limit of about 2.5 wt% Mn is preferred in the present invention. This upper limit is also preferred to substantially minimize centerline segregation that tends to occur in high Mn and continuously cast steels and the attendant through-thickness non-uniformity in microstructure and properties. More preferably, the upper limit for Mn content is about 2J wt%. If nickel content is increased above about 3 wt%, the desired high strength can be achieved without the addition of manganese. Therefore, in a broad sense, up to about 2.5 wt% manganese is preferred.

Silicon (Sϊ) may be added to steel for deoxidation purposes and a minimum of about 0.01 wt% is preferred for this purpose. However, Si is a strong BCC stabilizer and thus raises DBTT and also has an adverse effect on the toughness. For these reasons, when Si is added, an upper limit of about 0.5 wt% Si is preferred. More preferably, when Si is added, the upper limit for Si content is about 0J wt%. Silicon is not always necessary for deoxidation since aluminum or titanium can perform the same function. Niobium (Nb) is added to promote grain refinement of the rolled microstructure of the steel, which improves both the strength and toughness. Niobium carbide and carbonitride precipitation during hot rolling serves to retard recrystallization and to inhibit grain growth, thereby providing a means of austenite grain refinement. Also, precipitation of carbides and carbonitrides of niobium during tempering provides the desired secondary hardening to offset the strength loss normally observed in steel when it is tempered above about 500°C (930°F). For these reasons, at least about 0.02 wt% Nb is preferred, and at least about 0.03 wt% Nb is even more preferred. In some cases to obtain the desired strength after tempering, up to about 0.05 wt% Nb, and even up to about 0.08 wt% of NB, is preferred. However, Nb is a strong BCC stabilizer and thus raises DBTT. Too much Nb can be harmful to the weldability and HAZ toughness, so a maximum of about 0J wt% is preferred.

Vanadium (V is sometimes added to give precipitation strengthening by forming fine particles of the carbides and carbonitrides of vanadium in the steel on tempering and in its HAZ on cooling after welding. When dissolved in austenite, V has a strong beneficial effect on hardenability. When V is added to the steels of the present invention, at least about 0.02 wt% V is preferred. In some cases to obtain the desired strength after tempering, up to about 0.05 wt% V, and even up to about 0.08 wt% of V, is preferred. However, excessive V will help cause cold cracking on welding, and also deteriorate toughness of the base steel and its HAZ. The V addition, therefore, is preferably limited to a maximum of about 0J wt%.

Titanium (Ti). when added in a small amount, is effective in forming fine titanium nitride (TiN) particles which refme the grain size in both the rolled structure and the HAZ of the steel. Thus, the toughness of the steel is improved. Ti is added in such an amount that the weight ratio of Ti/N is preferably about 3.4. Ti is a strong BCC stabilizer and thus raises DBTT. Excessive Ti tends to deteriorate the toughness of the steel by forming coarser TiN or titanium carbide (TiC) particles. A Ti content below about 0.008 wt% generally can not provide sufficiently fine grain size or tie up the N in the steel as TiN while more than about 0.03 wt% can cause deterioration in toughness. More preferably, the steel contains at least about 0.01 wt% Ti and no more than about 0.02 wt% Ti.

Aluminum (AD is added to the steels of this invention for the purpose of deoxidation. At least about 0.001 wt% Al is preferred for this purpose, and at least about 0.005 wt% Al is even more preferred. Al also ties up nitrogen dissolved in the HAZ. However, Al is a strong BCC stabilizer and thus raises DBTT. If the Al content is too high, i.e., above about 0.05 wt%, there is a tendency to form aluminum oxide (Al₂O₃) type inclusions, which tend to be harmful to the toughness of the steel and its HAZ. Even more preferably, the upper limit for Al content is about 0.03 wt%. Molybdenum (Mo) increases the hardenability of steel on direct quenching, especially in combination with boron and niobium. Mo is also desirable for promoting secondary hardening during tempering of the steel by providing fine Mo C carbides. At least about 0J wt% Mo is preferred, and at least about 0J wt% Mo is even more preferred for this purpose. However, Mo is a strong BCC stabilizer and thus raises DBTT. Excessive Mo helps to cause cold cracking on welding, and also tends to deteriorate the toughness of the steel and HAZ, so a maximum of about 0.8 wt% is preferred, and a maximum of about 0.5 wt% is even more preferred. Therefore, in a broad sense, up to about 0.8 wt% Mo is preferred. Chromium (Cr) tends to increase the hardenability of steel on direct quenching. It also improves corrosion resistance and hydrogen induced cracking (HIC) resistance. Similar to Mo, excessive Cr tends to cause cold cracking in weldments, and also tends to deteriorate the toughness of the steel and its HAZ, so when Cr is added, a maximum of about 1.0 wt% Cr is preferred. More preferably, when Cr is added the Cr content is about 0J wt% to about 0.6 wt%.

Nickel (Ni is an important alloying addition to the steels of the present invention to obtain the desired DBTT, especially in the HAZ. It is one of the strongest FCC stabilizers in steel. Ni addition to the steel enhances the cross slip and thereby lowers DBTT. Although not to the same degree as Mn and Mo additions, Ni addition to the steel also promotes hardenability and therefore through-thickness uniformity in microstructure and properties in thick sections (i.e., thicker than about 25 mm (1 inch)). For achieving the desired DBTT in the weld HAZ, the minimum Ni content is preferably about 1.0 wt%, more preferably about 1.5 wt%. Since Ni is an expensive alloying element, the Ni content of the steel is preferably less than about 3.0 wt%, more preferably less than about 2.5 wt%, more preferably less than about 2.0 wt%, and even more preferably less than about 1.8 wt%, to substantially minimize cost of the steel.

Copper (Cu^') is a useful alloying addition to provide hardening during tempering via ε-copper precipitation. Preferably at least about 0J wt%, more preferably at least about 0.5 wt%, of Cu is added for this purpose. Cu is also an FCC stabilizer in steel and can contribute to lowering of DBTT in small amounts. Cu is also beneficial for corrosion and HIC resistance. At higher amounts, Cu induces excessive precipitation hardening and can lower the toughness and raise the DBTT both in the base plate and HAZ. Higher Cu can also cause embrittlement during slab casting and hot rolling, requiring co-additions of Ni for mitigation. For the above reasons, an upper limit of about 1.5 wt% Cu is preferred, and an upper limit of about 1.0 wt% is even more preferred.

Boron B) in small quantities can greatly increase the hardenability of steel and promote the formation of steel microstructures of lath martensite, lower bainite, and ferrite by suppressing the formation of upper bainite both in the base plate and the coarse grained HAZ. Generally, at least about 0.0004 wt% B is needed for this purpose. When boron is added to steels of this invention, from about 0.0006 wt% to about 0.0020 wt% is preferred, and an upper limit of about 0.0015 wt% is even more preferred. However, boron may not be a required addition if other alloying in the steel provides adequate hardenability and the desired microstructure.

DESCRIPTION AND EXAMPLES OF STEELS ACCORDING TO TfflS INVENTION

A 300 lb. heat of each chemical alloy shown in Table II was vacuum induction melted (VIM), cast into either round ingots or slabs of at least 130 mm thickness and subsequently forged or machined to 130 mm by 130 mm by 200 mm long slabs. The slabs were TMCP processed in a laboratory mill as described below. Table II shows the chemical composition of the alloys used for the TMCP processing.

TABLE II

Alloy

Cl C2

C (wt%) 0.052 0.049

Mn (wt%) 1.29 1.19

Ni (wt%) 2.02 3.02

Mo (wt%) 0.21 0.20

Cu (wt%) 0.82 0.81

Nb (wt%) 0.031 0.032

Si (wt%) 0.08 0.12

Ti (wt%) 0.013 0.012

Al (wt%) 0.017 0.013

B (ppm) 10 9

O(ppm) 17 76

S(ppm) 17 15

N(ppm) 27 18

P(ppm) 10 10

V (wt%) 0.033 0.033

Nc 2.76 2.79 The slabs were first reheated in a temperature range from about 1000°C to about 1080°C (1832°F to about 1976°F) for about 1 hour prior to the start of rolling according to the TMCP schedules shown in Table III:

TABLE III

Pass Thickness (mm) Temperature, °C

After Pass Cl C2

0 130 1071 1004

1 117 972 974

2 100 962 962

Delay, turn piece on the side

3 85 868 868

4 72 857 858

5 61 848 848

6 51 837 838

7 43 828 828

8 36 818 816

9 30 806 808

10 25 794 800

QST (°C) Ambient temperature

Cooling rate to QST (°C/s) 29 29

Tempering Temperature, °C 550 500

Following the preferred TMCP processing shown in Table III, the plate samples were tempered for 30 minutes at the indicated temperatures followed by quenching in water to ambient temperature. The microstructure of plate samples Cl and C2 is predominantly tempered FGB consistent with their low Nc. Cu-rich precipitates and cementite were the principal tempered precipitates in the FGB of both the tempered plates.

The transverse tensile strength and DBTT of the two plates of Tables II and III are summarized in Table IV. The tensile strengths and DBTTs summarized in Table IV were measured in the transverse direction, i.e., a direction that is in the plane of rolling but peφendicular to the plate rolling direction, wherein the long dimensions of the tensile test specimen and the Chaφy V-Notch test bar were substantially parallel to this direction with the crack propagation substantially peφendicular to this direction. A significant advantage of this invention is the ability to obtain the DBTT values summarized in Table IV in the transverse direction in the manner described in the preceding sentence. Consistent with the plates' low Nc and the consequent promotion of a predominant FGB microstructure in the plates, both the plates demonstrated strength at about and exceeding the minimum target of about 830 MPa (120 ksi) and at this strength good DBTT in the transverse direction is obtained for both plates although the DBTT of plate Cl is superior to that of plate C2. This is consistent with, and attributed in part, to the lower oxygen in the Cl plate and the significantly higher oxygen content of the C2 plate.

TABLE IV

Allov Cl C2

Tensile Strength. MPa (ksi) 847_. 854

023) 024)

DBTT. °C (°F) ______ -19 (-166) (-110)

Without thereby limiting this invention, the DBTT values given in TABLE IV correspond to the 50% energy transition temperature experimentally determined from Chaφy V-Notch impact testing according to standard procedures as set forth in

ASTM specification E-23, as will be familiar to those skilled in the art. The Chaφy V-Notch impact test is a well-known test for measuring the toughness of steels. This step-out combination of properties in the steels of the present invention provides a low cost enabling technology for certain cryogenic temperature operations, for example, storage and transport of natural gas at low temperatures. These new steels can provide significant material cost savings for cryogenic temperature applications over the current state-of-the-art commercial steels, which generally require far higher nickel contents (up to about 9 wt%) and are of much lower strengths (less than about 830 MPa (120 ksi)). Chemistry and microstructure design are used to lower DBTT and provide thick section capability for section thicknesses equal to or exceeding about 25 mm (1 inch). These new steels preferably have nickel contents lower than about 3 wt%, tensile strength greater than about 830 MPa (120 ksi), preferably greater than about 860 MPa (125 ksi), more preferably greater than about 900 MPa (130 ksi), and even more preferably greater than about 1000 MPa (145 ksi), ductile to brittle transition temperatures (DBTTs) for base metal in the transverse direction below about -62°C (-80°F), preferably below about -73°C

(-100°F), more preferably below about -100°C (-150°F), and even more preferably below about -123°C (-190°F), and offer excellent toughness at DBTT. Nickel content of these steel can be increased above about 3 wt% if desired to enhance performance after welding. Each 1 wt% addition of nickel is expected to lower the DBTT of the steel by about 10°C (18°F). Nickel content is preferably less than 9 wt%, more preferably less than about 6 wt%. Nickel content is preferably minimized in order to minimize cost of the steel.

While the foregoing invention has been described in terms of one or more preferred embodiments, it should be understood that other modifications may be made without departing from the scope of the invention, which is set forth in the following claims. Glossary of terms:

Ac^*ι transformation temperature: the temperature at which austenite begins to form during heating;

Ac₃ transformation temperature: the temperature at which transformation of ferrite to austenite is completed during heating;

AF: acicular ferrite;

Al₂O₃: aluminum oxide;

Ar transformation temperature: the temperature at which austenite begins to transform to ferrite during cooling;

BCC: body-centered cubic;

cementite: iron-rich carbide;

cooling rate: cooling rate at the center, or substantially at the center, of the plate thickness;

CRSS (critical resolved shear stress): an intrinsic property of a steel, sensitive to the ease with which dislocations can cross slip upon deformation, that is, a steel in which cross slip is easier will also have a low CRSS and hence a low DBTT;

cryogenic temperature: any temperature lower than about -40°C (-40°F); DBTT (Ductile to Brittle Transition Temperature): delineates the two fracture regimes in structural steels; at temperatures below the DBTT, failure tends to occur by low energy cleavage (brittle) fracture, while at temperatures above the DBTT, failure tends to occur by high energy ductile fracture;

DF: deformed ferrite;

DUB: deformed upper bainite;

effective grain size: as used in describing this invention, refers to mean austenite pancake thickness upon completion of rolling in the TMCP according to this invention and to mean packet width or grain size upon completion of transformation of the austenite pancakes to packets of lath martensite/lower bainite or FGB, respectively

FCC: face-centered cubic;

FGB (fine granular bainite): as used in describing this invention, an aggregate comprising bainitic ferrite as a major constituent and particles of mixtures of martensite and retained austenite as minor constituents;

gram: an individual crystal in a polycrystalline material; grain boundary: a narrow zone in a metal corresponding to the transition from one crystallographic orientation to another, thus separating one grain from another;

hardening particles one or more of ε-copper, Mo₂C, or the carbides and carbonitrides of niobium and vanadium;

HAZ: heat affected zone;

HIC: hydrogen induced cracking;

high angle boundary: a boundary that effectively behaves as a high angle grain boundary, i.e., a boundary that tends to deflect a propagating crack or fracture and, thus, induces tortuosity in a fracture path;

high angle grain boundary: a grain boundary that separates two adjacent grains whose crystallographic orientations differ by more than about 8°;

HSLA: high strength, low alloy;

mtercritically reheated: heated (or reheated) to a temperature of from about the Aci transformation temperature to about the Ac transformation temperature;

low alloy steel: a steel containing iron and less than about 10 wt% total alloy additives;

low heat input welding: welding with arc energies of up to about 2.5 kJ/mm (7.6 kJ/inch); MA: martensite-austenite;

major: as used in describing the present invention, means at least about 50 volume percent;

minor: as used in describing the present invention, means less than about 50 volume percent;

Mo₂C: a form of molybdenum carbide;

M_s transformation temperature: the temperature at which transformation of austenite to martensite starts during cooling;

Nc: a factor defined by the chemistry of the steel as {Nc = 12.0*C + Mn + 0.8*Cr + 0.15*(Ni + Cu)+ 0.4*Si + 2.0*V + 0.7* b + 1.5*Mo}, where C, Mn, Cr, Ni, Cu, Si, V, Nb, Mo represent their respective wt % in the steel;

PF: polygonal ferrite;

predominantly/predominant: as used in describing the present invention, means at least about 50 volume percent;

prior austenite grain size: average austenite grain size in a hot-rolled steel plate prior to rolling in the temperature range in which austenite does not recrystallize; quenching: as used in describing the present invention, accelerated cooling by any means whereby a fluid selected for its tendency to increase the cooling rate of the steel is utilized, as opposed to air cooling;

Quench Stop Temperature (QST): the highest, or substantially the highest, temperature reached at the surface of the plate, after quenching is stopped, because of heat transmitted from the mid-thickness of the plate;

slab: a piece of steel having any dimensions;

Sv : total interfacial area of the high angle boundaries per unit volume in steel plate;

tensile strength: in tensile testing, the ratio of maximum load to original cross-sectional area;

thick section capability: the ability to provide substantially the desired microstructure and properties (e.g., strength and toughness), particularly in thicknesses equal to or greater than about 25 mm (1 inch);

through-thickness direction: a direction that is orthogonal to the plane of rolling;

TiC: titanium carbide;

TiN: titanium nitride; Tn_j- temperature: the temperature below which austenite does not recrystallize;

TMCP: thermo-mechanical controlled rolling processing;

transverse direction: a direction that is in the plane of rolling but peφendicular to the plate rolling direction;

UB: upper bainite; and

VIM: vacuum induction melted.

Claims

We Claim:

1. A method for preparing a steel plate having a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered fine granular bainite (FGB), or mixtures thereof, said method comprising the steps of:

(a) heating a steel slab to a reheating temperature sufficiently high to (i) substantially homogenize said steel slab, (ii) dissolve substantially all carbides and carbonitrides of niobium and vanadium in said steel slab, and (iii) establish fine initial austenite grains in said steel slab;

(b) reducing said steel slab to form steel plate in one or more hot rolling passes in a first temperature range in which austenite recrystallizes;

(c) further reducing said steel plate in one or more hot rolling passes in a second temperature range below about the Tn_j- temperature and above about the Ar₃ transformation temperature;

(d) quenching said steel plate at a cooling rate of at least about 10°C per second (18°F/sec) to a Quench Stop Temperature below about 600°C

(1112°F);

(e) stopping said quenching; and

(f) tempering said steel plate at a tempering temperature from about 400°C (752°F) to about the Aci transformation temperature for a period of time sufficient to cause precipitation of hardening particles, said steps being performed so as to facilitate transformation of said microstructure of said steel plate to predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered fine granular bainite (FGB), or mixtures thereof.

2. The method of claim 1 wherein step (f) is replaced with the following:

(f) tempering said steel plate at a tempering temperature from about 400°C (752°F) to about the Aci transformation temperature for a period of time sufficient to cause precipitation of hardening particles, said steps being performed so as to facilitate transformation of said microstructure of said steel plate to substantially 100% tempered fine-grained lath martensite.

3. The method of claim 1 wherein step (f) is replaced with the following:

(f) tempering said steel plate at a tempering temperature from about

400°C (752°F) to about the Ac] transformation temperature for a period of time sufficient to cause precipitation of hardening particles, said steps being performed so as to facilitate transformation of said microstructure of said steel plate to predominantly fine granular bainite (FGB).

4. The method of claim 1 wherein said reheating temperature of step (a) is between about 955°C and about 1100°C (1750°F - 2012°F).

5. The method of claim 1 wherein said fine initial austenite grains of step (a) have a grain size of less than about 120 microns.

6. The method of claim 1 wherein a reduction in thickness of said steel slab of about 30% to about 70% occurs in step (b).

7. The method of claim 1 wherein a reduction in thickness of said steel plate of about 40%) to about 80% occurs in step (c).

8. The method of claim 1 further comprising the step of allowing said steel plate to air cool to ambient temperature from said Quench Stop Temperature prior to tempering said steel plate in step (f).

9. The method of claim 1 wherein said steel slab of step (a) comprises iron and the following alloying elements in the weight percents indicated: about 0.03% to about 0.12% C, at least about 1% Ni, up to about 1.5% Cu, up to about 0.8% Mo, about 0.02% to about 0.1% Nb, about 0.008% to about 0.03% Ti, about 0.001% to about 0.05% Al, and about 0.002% to about 0.005% N.

10. The method of claim 9 wherein said steel slab comprises less than about 6 wt% Ni.

11. The method of claim 9 wherein said steel slab comprises less than about 3 wt% Ni and additionally comprises about 0.5 wt% to about 2.5 wt% Mn.

12. The method of claim 9 wherein said steel slab further comprises at least one additive selected from the group consisting of (i) up to about 1.0 wt% Cr, (ii) up to about 0.5 wt% Si, (iii) up to about 0.1 wt% V, (iv) up to about 2.5 wt% Mn, and about 0.0004 wt% to about 0.0020 wt% B.

13. The method of claim 9 wherein said steel slab further comprises about 0.0004 wt% to about 0.0020 wt% B.

14. The method of claim 1 wherein, after the tempering of step (f), said steel plate has a DBTT lower than about -62°C (-80°F) in both said steel plate and its HAZ and has a tensile strength greater than about 830 MPa (120 ksi).

15. A steel plate having a DBTT of lower than about -62°C (-80°F) in both said steel plate and its HAZ, a tensile strength greater than about 830 MPa (120 ksi), and a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, tempered fine granular bainite (FGB), or mixtures thereof, and wherein said steel plate is produced from a reheated steel slab comprising iron and the following alloying elements in the weight percents indicated: about 0.03% to about 0.12% C, at least about 1% Ni, up to about 1.5%) Cu, up to about 0.8% Mo, about 0.02% to about 0.1% Nb, about 0.008% to about 0.03% Ti, about 0.001% to about 0.05% Al, and about 0.002% to about 0.005% N.

16. The steel plate of claim 15 wherein said steel slab comprises less than about 6 wt% Ni.

17. The steel plate of claim 15 wherein said steel slab comprises less than about 3 wt% Ni and additionally comprises about 0.5 wt% to about 2.5 wt% Mn.

18. The steel plate of claim 15 further comprising at least one additive selected from the group consisting of (i) up to about 1.0 wt% Cr, (ii) up to about 0.5 wt% Si, (iii) up to about 0.1 wt% V, (iv) up to about 2.5 wt% Mn, and (v) about 0.0004 wt% to about 0.0020 wt% B.

19. The steel plate of claim 15 further comprising about 0.0004 wt% to about 0.0020 wt% B.

20. A method for obtaining a DBTT of lower than about -62°C (-80°F) in the HAZ of steel plate by adding at least about 1.0 wt% Ni and at least about 0.1 wt% Cu, and by substantially minimizing addition of BCC stabilizing elements.

21. A method for preparing a steel plate having a microstructure comprising predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, or mixtures thereof, said method comprising the steps of:

(a) heating a steel slab to a reheating temperature sufficiently high to (i) substantially homogenize said steel slab, (ii) dissolve substantially all carbides and carbonitrides of niobium and vanadium in said steel slab, and (iii) establish fine initial austenite grains in said steel slab;

(b) reducing said steel slab to form steel plate in one or more hot rolling passes in a first temperature range in which austenite recrystallizes;

(c) further reducing said steel plate in one or more hot rolling passes in a second temperature range below about the T^,- temperature and above about the Ar₃ transformation temperature;

(d) quenching said steel plate at a cooling rate of about 10°C per second to about 40°C per second (18°F/sec - 72°F/sec) to a Quench Stop Temperature below about the M_s transformation temperature plus 200°C (360°F);

(e) stopping said quenching; and (f) tempering said steel plate at a tempering temperature from about 400°C (752°F) to about the Aci transformation temperature for a period of time sufficient to cause precipitation of hardening particles, said steps being performed so as to facilitate transformation of said microstructure of said steel plate to predominantly tempered fine-grained lath martensite, tempered fine-grained lower bainite, or mixtures thereof.

22. A method for controlling the mean ratio of pancake length to pancake thickness during processing of a an ultra-high strength, steel plate in order to enhance transverse toughness and transverse DBTT of said steel plate, said method comprising the steps of:

(a) heating a steel slab to a reheating temperature sufficiently high to (i) substantially homogenize said steel slab, (ii) dissolve substantially all carbides and carbonitrides of niobium and vanadium in said steel slab, and (iii) establish fine initial austenite grains in said steel slab;

(b) reducing said steel slab to form steel plate in one or more hot rolling passes in a first temperature range in which austenite recrystallizes;

(c) further reducing said steel plate in one or more hot rolling passes in a second temperature range below about the T^- temperature and above about the Ar₃ transformation temperature;

(d) quenching said steel plate at a cooling rate of at least about 10°C per second (18°F/sec) to a Quench Stop Temperature below about 600°C (1112°F);

(e) stopping said quenching, so as to produce a mean ratio of pancake length to pancake thickness of less than about 100 in said steel plate.