Patent Publication Number: US-2017349983-A1

Title: High strength cryogenic high manganese steels and methods of making the same

Description:
CROSS-REFERENCE TO RELATED APPLICATIONS 
     This application claims the benefit of U.S. Provisional Application No. 62/346,069, filed on Jun. 6, 2016, the entire contents of which is incorporated herein by reference. 
    
    
     FIELD 
     The present disclosure relates to cost-effective high manganese steels tailored to achieve high yield strength, cryogenic toughness, and low thermal stress. More specifically, the present disclosure pertains to ferrous steel alloyed with a high amount (≧8 wt. %) of manganese providing high yield strength, cryogenic toughness, and low thermal stress and manufacturing the same. The strength, toughness, elastic modulus, thermal expansion coefficient and thermal conductivity of these steels can be optimized through the control of microstructure and chemistry. The application of the present disclosure includes, but is not limited to liquefied natural gas (LNG) storage, piping, and transfer by pipeline. 
     BACKGROUND 
     Cryogenic structures such as liquefied natural gas (LNG) container vessels demand steels with specific low temperature properties. The steels need to remain ductile and crack resistant and must retain a high level of safety even at cryogenic temperatures (&lt;−50° C.). These steels must also possess high strength in order to allow reduction of tank wall thickness which permits low cost construction. Conventional carbon steels lose much of their toughness and become brittle at cryogenic temperatures. Steels commonly used for structural applications at cryogenic temperatures include alloy steels such as Fe-9 wt. % Ni steel, austenitic stainless steels (e.g., 304 SS with Fe-18 wt. % Cr-8 wt. % Ni), invar alloys (Fe-36 wt % Ni), and aluminum alloys. 
     Aluminum alloys are used in various cryogenic applications due to their high specific strength and ductility. However, most aluminum alloys have lower strength compared with the strength of alloyed steel and are relatively challenging to weld. Austenitic stainless steels (e.g., 304 SS) and Invar alloys are relatively low strength and high cost. Nickel-alloyed high-strength steels (5% Ni and 9% Ni) provide a combination of high cryogenic strength and toughness and, hence, 9% Ni steels are often preferred for the most demanding low temperature applications. However, because of high Ni content, these alloys are expensive. 
     Ferrous steels with high manganese (Mn) alloying can be a lower-cost alternative to these cryogenic materials. The essential benefits of the inventive steels of the present disclosure are the lower cost by replacing Ni with Mn, and higher strength due to higher carbon (C), and/or nitrogen (N) contents than the steels in the art. 
     Liquefied natural gas is usually transported by specially equipped ships and stored in LNG-terminals. Conventional LNG carrier ships are of two basic types. The first type (Moss type) uses heavy walled self-supporting spherical tanks made of aluminum to contain the LNG. The second type (membrane type) uses a thin membrane of Invar alloy or corrugated austenitic stainless steel supported by the ship&#39;s hull (with plywood and insulation in between) to contain the LNG. For export and import terminals, free standing LNG storage tanks are typically made out of 9% Ni steel. In LNG tanks, cryogenic steels are joined by fusion welding technology for liquid-tightness. Typical joining techniques utilized in the field are GTAW (Gas Tungsten Arc Welding), GMAW (Gas Metal Arc Welding), SMAW (Shielded Metal Arc Welding), and SAW (Submerged Arc Welding). 
     Even though extensive studies have been made on welding technologies for cryogenic steels, it remains challenging to cost effectively meet weld property requirements in cryogenic steel weldments. In the case of 9% Ni steel, for instance, achieving stable cryogenic toughness in the as-welded state (without heat treatment) can be challenging when a weldment is fabricated with similar composition filler wire. For this reason, Ni-based alloy composition weld wire is typically used for joining 9% Ni steels. Weldments with Ni-based alloy weld wire, however, show lower yield strength than that of 9% Ni steel itself, thus compromising the full strength of the 9% Ni steel. Furthermore, weldments with Ni-based weld wires can be susceptible to high temperature cracking (during welding) and fatigue damage due to a difference in thermal expansion coefficients. In addition, high nickel content increases the cost of welding consumables. 
     Traditional LNG loading/offloading lines of stainless steel require mechanical expansion loops and bellows that can deflect with thermal stress to accommodate thermal contraction/expansion upon temperature excursion between construction and operating temperatures (refer to  FIG. 1 ). For LNG loading terminals, thus, a driving force exists to change loading pipeline design from the standard jetty-based stainless steel to a system that does not require such accommodations (i.e., mechanical expansion loops and bellows). 
     By switching to alloys with higher yield strength and lower thermal expansion coefficients, these expansion loops can be eliminated along with the above-sea trestle required to accommodate them. The lower coefficient of thermal expansion (CTE) value decreases thermal stresses that arise when the pipeline is cooled from ambient to operating temperature. Developing such pipelines with 9% Ni steel has been attempted, but as for 9% Ni tanks designs, problems are incurred due to the necessity of welding 9% Ni steel with undermatching, austenitic welding consumables. The inventive concept of using high Mn steels (which can be welded with matching or overmatching weld consuambles) for cryogenic pipeline design allows for numerous advantages including reduced materials and construction cost, increased pipeline security, reduced environmental impact, capital savings associated with jetty construction, and reduced operating costs. 
     SUMMARY 
     The present disclosure provides for a method for fabricating a ferrous based component comprising: a) providing a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition or to a temperature to homogenize the composition; c) cooling the composition to a rolling start temperature; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     The present disclosure also provides for a ferrous based component fabricated according to the steps comprising: a) providing a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
    
    
     
       BRIEF DESCRIPTION OF THE DRAWINGS 
       Features and aspects of embodiments are described below with reference to the accompanying drawings, in which elements are not necessarily depicted to scale. 
       Exemplary embodiments of the present disclosure are further described with reference to the appended figures. It is to be noted that the various steps, features and combinations of steps/features described below and illustrated in the figures can be arranged and organized differently to result in embodiments which are still within the spirit and scope of the present disclosure. To assist those of ordinary skill in the art in making and using the disclosed systems, assemblies and methods, reference is made to the appended figures, wherein: 
         FIG. 1  is a diagram of a LNG loading line with expansion loop, as required using “traditional” (i.e., NOT the inventive high Mn steel of the present disclosure) piping technologies and/or piping compositions; 
         FIG. 2  is an exemplary diagram of the phase stability and deformation mechanism of high Mn steels as a function of alloy chemistry and temperature; 
         FIG. 3  displays the predicted influence of alloying elements on the SFE values of the FeMn13C0.6 reference and the deformation mechanism; 
         FIG. 4  depicts the effect of carbide precipitates on mechanical properties and deformation mechanism (schematic, not in scale); 
         FIG. 5  is a schematic drawing of an exemplary steel fabrication method of the present disclosure; 
         FIG. 6  depicts a schematic of an exemplary fabrication method for producing ultrafine grained high Mn steels according to the present disclosure. 
     
    
    
     DETAILED DESCRIPTION 
     All numerical values within the detailed description and the claims herein are modified by “about” or “approximately” the indicated value, and take into account experimental error and variations that would be expected by a person having ordinary skill in the art. 
     Where a range of values is provided, it is understood that each intervening value, to the tenth of the unit of the lower limit unless the context clearly dictates otherwise, between the upper and lower limit of that range and any other stated or intervening value in that stated range is encompassed within the disclosure. Ranges from any lower limit to any upper limit are contemplated. The upper and lower limits of these smaller ranges which may independently be included in the smaller ranges is also encompassed within the disclosure, subject to any specifically excluded limit in the stated range. Where the stated range includes one or both of the limits, ranges excluding either both of those included limits are also included in the disclosure. 
     Although any methods and materials similar or equivalent to those described herein can also be used in the practice or testing of the present disclosure, the preferred methods and materials are now described. All publications mentioned herein are incorporated herein by reference to disclose and described the methods and/or materials in connection with which the publications are cited. 
     It must be noted that as used herein and in the appended claims, the singular forms “a”, “and”, and “the” include plural references unless the context clearly dictates otherwise. 
     Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this disclosure belongs. The terminology used in the description of the disclosure herein is for describing particular embodiments only and is not intended to be limiting of the disclosure. All publications, patent applications, patents, figures and other references mentioned herein are expressly incorporated by reference in their entirety. 
     Definitions 
     CRA: Corrosion resistant alloys, can mean, but is in no way limited to, a specially formulated material used for completion components likely to present corrosion problems. Corrosion-resistant alloys may be formulated for a wide range of aggressive conditions. 
     Ductility: can mean, but is in no way limited to, a measure of a material&#39;s ability to undergo appreciable plastic deformation before fracture; it may be expressed as percent elongation (% EL) or percent area reduction (% AR). 
     Erosion resistance: can mean, but is in no way limited to, a material&#39;s inherent resistance to erosion when exposed to moving solid particulates striking the surface of the material. 
     Toughness: can mean, but is in no way limited to, resistance to fracture initiation. 
     Fatigue: can mean, but is in no way limited to, resistance to fracture under cyclic loading. 
     Yield Strength: can mean, but is in no way limited to, the ability to bear load without deformation. 
     Cooling rate: can mean, but is in no way limited to, the rate of cooling at the center, or substantially at the center, of the plate thickness. 
     Austenite: can mean, but is in no way limited to, a solid solution of one or more elements in face-centered cubic crystallographic structure of iron; the solute can be, but not limited to, carbon, nitrogen, manganese, and nickel. 
     Martensite: can mean, but is in no way limited to, a generic term for microstructures formed by diffusionless phase transformation in which the parent (typically austenite) and product phases have a specific orientation relationship. 
     ε (epsilon)-martensite: can mean, but is in no way limited to, a specific form of martensite having hexagonal close packed crystal structure which forms upon cooling or straining of austenite phase. ε-martensite typically forms on close packed (111) planes of austenite phase and is similar to deformation twins or stacking fault clusters in morphology. 
     α′(alpha prime)-martensite: can mean, but is in no way limited to, a specific form of martensite having body centered cubic or body centered tetragonal crystal structure which forms upon cooling or straining of austenite phase; α′-martensite typically forms as platelets. 
     M s  temperature: can mean, but is in no way limited to, the temperature at which transformation of austenite to martensite starts during cooling. 
     M f  temperature: can mean, but is in no way limited to, the temperature at which transformation of austenite to martensite finishes during cooling. 
     M d  temperature: can mean, but is in no way limited to, the highest temperature at which a designated amount of martensite forms under defined deformation conditions. M d  temperature is typically used to characterize the austenite phase stability upon deformation. 
     Carbide: can mean, but is in no way limited to, a compound of iron/metal and carbon. 
     Cementite: can mean, but is in no way limited to, a compound of iron and carbon having approximate chemical formula of Fe 3 C with orthorhombic crystal structure. 
     Pearlite: can mean, but is in no way limited to, typically a lamellar mixture of two-phases, made up of alternate layers of ferrite and cementite (Fe 3 C). 
     Grain: can mean, but is in no way limited to, an individual crystal in a polycrystalline material. 
     Grain boundary: can mean, but is in no way limited to, a narrow zone in a metal corresponding to the transition from one crystallographic orientation to another, thus separating one grain from another. 
     Quenching: can mean, but is in no way limited 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. 
     Accelerated cooling start temperature (ACST): can mean, but is in no way limited to, the temperature reached at the surface of plate, when quenching is initiated. 
     Accelerated cooling finish temperature (ACFT): can mean, but is in no way limited to, 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. 
     Recrystallization: the formation of a new, strain-free grain structure grains from cold-worked metal accomplish by heating through a critical temperature. 
     T nr  temperature: the temperature below which austenite does not recrystallize. 
     The present disclosure provides advantageous steel compositions. More particularly, the present disclosure provides improved high manganese (Mn) steel having high yield strength/cryogenic toughness/low thermal stress, and related methods for fabricating steel having enhanced yield strength/cryogenic toughness/low thermal stress. In exemplary embodiments, the advantageous steel compositions/components of the present disclosure improve one or more of the following properties: yield strength, cryogenic toughness, thermal stress, elastic modulus, thermal expansion coefficient and/or thermal conductivity. 
     In general, the present disclosure provides for cost-effective high manganese steels having improved performance during cryogenic applications (e.g., yield strength, cryogenic toughness, thermal stress, elastic modulus, thermal expansion coefficient and/or thermal conductivity). More specifically, the present disclosure provides ferrous steel alloyed with a high amount (e.g., greater than or equal to about 5 weight %) of manganese, and where the fabricated steel exhibits increased/improved yield strength/cryogenic toughness/low thermal stress (e.g., improved yield strength, cryogenic toughness, thermal stress, elastic modulus, thermal expansion coefficient and/or thermal conductivity). The present disclosure also provides methods for fabricating such improved steel. In exemplary embodiments, the high manganese steels of the present disclosure have advantages/potential in applications where strength/toughness is desired/required at cryogenic temperatures. 
     In certain aspects, the disclosure provides methods for improving the yield strength, cryogenic toughness, and thermal stress of the steels through the control of microstructure and/or chemistry. In certain embodiments, the methods include steps to promote phase transformations (e.g., to alpha prime martensite or epsilon martensite phases), twinning during deformation, and/or introducing hard erosion resistant second phase particles to the compositions. 
     Some exemplary uses/applications of the steel compositions of present disclosure include, without limitation, are use in piping systems, material conveying systems, fluids/solids transport systems, in mining operations, and/or as material for liquefied natural gas (LNG) storage systems, and piping for LNG onloading/offloading piping systems. In another aspect, the liquefied natural gas (LNG) storage systems, and piping for LNG onloading/offloading piping systems are waffle-free, wrinkle-free, and/or free of expansion loop piping. Moreover, the use of the steels of the present disclosure can improve the economics of LNG storage and/or transport. 
     Exemplary Methods for Fabrication 
     The present disclosure provides for a method for fabricating a ferrous based component including: a) providing a composition having from about 5 to about 40 weight % manganese, preferably from about 9 to about 25 weight % manganese, even more preferably from about 12 to about 20 weight % manganese, and from about 0.01 to about 1.2 weight % carbon, preferably from about 0.3% to about 1.2 weight % carbon, even more preferably from about 0.3% to about 0.7 weight % carbon, and the balance iron, b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition (e.g., to a temperature to homogenize the composition); c) cooling to a rolling start temperature (RST) d) deforming or hot rolling the composition while the composition is at a temperature above the austenite recrystallization stop temperature of the composition; and e) quenching or accelerated cooling the composition. 
     The present disclosure provides for a method for fabricating a ferrous based component including: a) providing a composition having from about 5 to about 40 weight % manganese, preferably from about 9 to about 25 weight % manganese, even more preferably from about 12 to about 20 weight % manganese, and from about 0.01 to about 1.2 weight % carbon, preferably from about 0.3% to about 1.2 weight % carbon, even more preferably from about 0.3% to about 0.7 weight % carbon, and the balance iron, b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition (e.g., to a temperature to homogenize the composition); c) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and d) quenching or accelerated cooling the composition. In another aspect, the temperature recited in step c) is in the range of 700-1100° C., preferably about 800-1000° C., and more preferably about 800-900° C. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein after step d), the matrix of the composition is predominantly or substantially in the austenitic phase. In one or more embodiments, the volume percent of austenite in the steel composition is from about 50 volume % to about 100 volume %, more preferably from about 80 volume % to about 99 volume %, even more preferably from about 90 volume % to about 98 volume %. 
     The steel composition is preferably processed into predominantly or substantially austenitic plates using a hot rolling process. In one or more embodiments, a steel billet/slab from the compositions described is first formed, such as, for example, through a continuous casting process. The billet/slab can then be re-heated to a temperature (“reheat temperature”) within the range of about 1,000° C. to about 1,300° C., more preferably within the range of about 1050° C. to 1250° C., even more preferably within the range of about 1100° C. to 1200° C. Preferably, the reheat temperature is sufficient to: (i) substantially homogenize the steel slab/composition, (ii) dissolve substantially all the carbide and/or carbonitrides, when present, in the steel slab/composition, and (iii) establish fine initial austenite grains in the steel slab/composition. 
     The re-heated slab/composition can then be hot rolled in one or more passes. In exemplary embodiments, the rolling or hot deformation can be initiated at a “rolling start temperature”. In one or more embodiments, the rolling start temperature is above 1100° C., preferably above 1080° C., even more preferably above 1050° C. In exemplary embodiments, the final rolling for plate thickness reduction can be completed at a “rolling finish temperature”. In one or more embodiments, the rolling finish temperature is above about 700° C., preferably above about 800° C., more preferably about 850° C. Thereafter, the hot rolled plate can be cooled (e.g., in air) to a first cooling temperature or accelerated cooling start temperature (“ACST”), at which an accelerated cooling starts to cool the plates at a rate of at least about 10° C. per second to a second cooling temperature or accelerated cooling finish temperature (“ACFT”). After the cooling to the ACFT, the steel plate/composition can be cooled to room temperature (e.g., ambient temperature) in ambient air. Preferably, the steel plate/composition is allowed to cool on its own to room temperature. 
     In one or more embodiments, the ACST is about 750° C. or more, about 800° C. or more, about 850° C. or more, or about 900° C. or more. In one or more embodiments, the ACST can range from about 700° C. to about 1000° C. In one or more embodiments, the ACST can range from about 750° C. to about 950° C. Preferably, the ACST ranges from a low of about 650° C., 700° C., or 750° C. to a high of about 900° C., 950° C., or 1000° C. In one or more embodiments, the ACST can be about 750° C., about 800° C., about 850° C., about 890° C., about 900° C., about 930° C., about 950° C., about 960° C., about 970° C., about 980° C., or about 990° C. 
     In one or more embodiments, the ACFT can range from about 0° C. to about 500° C. Preferably, the ACFT ranges from a low of about 0° C., 10° C., or 20° C. to a high of about 150° C., 200° C., or 300° C. 
     Without being bound by any theory, it is believed that the rapid cooling (e.g., more than about 10° C./sec cooling rate) to the low accelerated cooling finish temperature (“ACFT”) retards at least a portion of the carbon and/or nitrogen atoms from diffusing from the austenite phase of the steel composition to the grain boundary or second phase. It is further believed that the high accelerated cooling start temperature (“ACST”) retards at least a portion of the carbon and/or nitrogen atoms from forming precipitates such as, for example, carbides, carbonitrides, and/or nitrides during subsequent cooling to the ACFT. As such, the amount of precipitates at the grain boundaries is reduced. Therefore, the steel&#39;s fracture toughness and/or resistance to cracking and/or strength at cryogenic temperatures are enhanced. 
     Following the rolling and cooling steps, the plate can be formed into pipes or the like (e.g., linepipe). Any suitable method for forming pipe can be used. Preferably, the precursor steel plate is fabricated into linepipe by a conventional UOE process or JCOE process which is known in the art. 
     The present disclosure also provides for a method for fabricating a ferrous based component including: a) providing a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     The present disclosure also provides for a method for fabricating a ferrous based component wherein after step e), the carbide precipitate fraction volume of the composition is about 3 volume % or less of the composition, preferably about 1.2 volume % or less of the composition, and even more preferably about 0.7 volume % or less of the composition. The present disclosure also provides for a method for fabricating a ferrous based component wherein after step e), the composition has a microstructure having a refined grain size of about 100 μm or less, preferably about μm or less, even more preferably about 30 μm or less. 
     The present disclosure also provides for a method for fabricating a ferrous based component wherein the microstructure having a refined grain size of about 100 μm or less includes a surface layer of the composition. The present disclosure also provides for a method for fabricating a ferrous based component wherein the thickness of the surface layer is from about 10 nm to about 5000 nm. The present disclosure also provides for a method for fabricating a ferrous based component wherein the surface layer is formed prior to or during the use of the composition. The present disclosure provides for a method for fabricating a ferrous based component wherein the surface layer is formed via a surface deformation technique selected from the group consisting of shot peening, laser shock peening, surface burnishing and combinations thereof. The present disclosure provides for a method for fabricating a ferrous based component further including after step e) a surface deformation step selected from the group consisting of shot peening, laser shock peening, surface burnishing and combinations thereof. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein prior to step e), the composition is slowly cooled or isothermally held. The present disclosure provides for a method for fabricating a ferrous based component wherein step e) includes rapidly quenching the composition. The present disclosure provides for a method for fabricating a ferrous based component wherein step d) includes deforming the composition while the composition is at a temperature below the austenite recrystallization temperature and above the martensite transformation start temperature. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein step d) includes deforming the composition to induce martensite formation of the composition. The present disclosure provides for a method for fabricating a ferrous based component wherein the composition is deformed at a temperature of from about 18° C. to about 24° C. to induce martensite formation of the composition. The present disclosure provides for a method for fabricating a ferrous based component further including, after step d), heating the composition to a temperature above the austenite recrystallization stop temperature. The present disclosure provides for a method for fabricating a ferrous based component wherein heating the composition to a temperature above the austenite recrystallization stop temperature after step d) reverses deformation-induced martensite of the composition into ultrafine grained austenite. The present disclosure provides for a method for fabricating a ferrous based component wherein the martensite start temperature of the ultrafine grained austenite is below about 24° C. 
     The present disclosure provides for a method for fabricating a ferrous based component further including, after step e), heating the composition to a temperature above the austenite recrystallization stop temperature, and then quenching the composition. The present disclosure provides for a method for fabricating a ferrous based component further including, prior to step c), deforming the composition while the composition is at a temperature above the austenite recrystallization stop temperature. The present disclosure provides for a method for fabricating a ferrous based component wherein the composition is deformed at a temperature of from about 700° C. to about 1000° C. The present disclosure provides for a method for fabricating a ferrous based component wherein step b) includes heating the composition to at least about 1000° C. The present disclosure provides for a method for fabricating a ferrous based component wherein step c) includes cooling the composition at a rate of from about 2° C. per second to about 60° C. per second. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein the composition further includes one or more alloying elements selected from the group consisting of chromium, aluminum, silicon, nickel, cobalt, molybdenum, niobium, copper, titanium, vanadium, nitrogen, boron, zirconium, hafnium and combinations thereof. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein the chromium ranges from 0 to 30 weight % of the total composition, more preferably from 0.5 to 20 weight % of the total composition, even more preferably from 2 to 5 weight % of the total composition; wherein each of the nickel or cobalt ranges from 0 to 20 weight % of the total composition, more preferably from 0.5 to 20 weight % of the total composition, even more preferably from 1 to 5 weight % of the total composition; wherein the aluminum ranges from 0 to 15 weight % of the total composition, more preferably from 0.5 to 10 weight % of the total composition, even more preferably from 1 to 5 weight % of the total composition; wherein each of the molybdenum, niobium, copper, titanium or vanadium ranges from 0 to 10 weight % of the total composition, more preferably from 0.02 to 5 weight % of the total composition, even more preferably from 0.02 to 2 weight % of the total composition; wherein the silicon ranges from 0 to 10 weight % of the total composition, more preferably from 0.1 to 6 weight % of the total composition, even more preferably from 0.1 to 0.5 weight % of the total composition; wherein the nitrogen ranges from 0 to 3.0 weight % of the total composition, more preferably from 0.01 to 2.0 weight % of the total composition, even more preferably from 0.01 to 1.0 weight % of the total composition; wherein the boron ranges from 0 to 0.1 weight % of the total composition, more preferably from 0.001 to 0.1 weight % of the total composition; and wherein each of the zirconium or hafnium ranges from 0 to 6 weight % (e.g., 0.2 to 5 wt %) of the total composition. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein the composition includes from about 8 to about 20 weight % manganese, from about 0.30 to about 0.7 weight % carbon, from about 0.5 to about 5 weight % chromium, from about 0.0 to about 2.0 weight % copper, from about 0.01 to about 1 weight % silicon, and the balance iron. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein step d) includes transformation induced plasticity or twin-induced plasticity. 
     The present disclosure also provides for a ferrous based component including a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; and wherein the carbide precipitate fraction volume of the composition is about 2 volume % or less of the composition. 
     The present disclosure also provides for a ferrous based component including a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; and wherein the composition has a microstructure having a refined grain size of about 30 μm or less. 
     The present disclosure also provides for a ferrous based component fabricated according to the steps comprising: a) providing a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     Any combination or permutation of embodiments is envisioned. Additional advantageous features, functions and applications of the disclosed systems and methods of the present disclosure will be apparent from the description which follows, particularly when read in conjunction with the appended figures. All references listed in this disclosure are hereby incorporated by reference in their entireties. 
     Exemplary Embodiments 
     The exemplary embodiments disclosed herein are illustrative of advantageous steel compositions, and systems of the present disclosure and methods/techniques thereof. It should be understood, however, that the disclosed embodiments are merely exemplary of the present disclosure, which may be embodied in various forms. Therefore, details disclosed herein with reference to exemplary steel compositions/fabrication methods and associated processes/techniques of assembly and use are not to be interpreted as limiting, but merely as the basis for teaching one skilled in the art how to make and use the advantageous steel compositions of the present disclosure. Drawing figures are not necessarily to scale and in certain views, parts may have been exaggerated for purposes of clarity. 
     The present disclosure provides advantageous steel compositions (e.g., having enhanced strength/performance at cryogenic temperatures). More particularly, the present disclosure provides improved high manganese (Mn) steel having high yield strength/cryogenic toughness/low thermal stress, and methods for fabricating high manganese steel compositions having high yield strength/cryogenic toughness/low thermal stress. In exemplary embodiments, the advantageous steel compositions/components of the present disclosure improve one or more of the following properties: strength, toughness, elastic modulus, thermal expansion coefficient and/or thermal conductivity. 
     Representative properties of commercial cryogenic materials and inventive high Mn austenitic steels (i.e., high Mn steels of the present disclosure) are shown in Table 1. 
     
       
         
           
               
             
               
                 TABLE 1 
               
             
            
               
                   
               
               
                 Typical properties of cryogenic materials and inventive steels. 
               
            
           
           
               
               
               
               
               
               
            
               
                   
                   
                   
                   
                 Cryogenic 
                 Inventive 
               
               
                   
                   
                   
                 9% Ni 
                 High Mn 
                 High Mn 
               
               
                 Properties 
                 304L SS 
                 Al5083 
                 steel 
                 steel 
                 Steel 
               
               
                   
               
            
           
           
               
               
               
               
               
               
            
               
                 Yield Strength (MPa) 
                 ≧190 
                 ≧142 
                 ≧585 
                 ≧400 
                 ≧600 
               
               
                 Ultimate Tensile 
                 ≧420 
                 ≧290 
                 ≧690 
                 ≧560 
                 ≧1050 
               
               
                 Strength (MPa) 
               
               
                 Charpy toughness (J@ 
                 ≧60 
                 ≧30 
                 ≧100 
                 ≧82 
                 ≧110 (@ 
               
               
                 −196° C.) 
                   
                   
                   
                   
                 −40° C.) 
               
               
                   
                   
                   
                   
                   
                 ≧56 (@ 
               
               
                   
                   
                   
                   
                   
                 −196° C.) 
               
               
                 Elastic modulus (GPa) 
                 ~193 
                 ~70 
                 ~186 
                 ~206 
                 ~190 
               
               
                 Thermal Expansion 
                 ~15 
                 ~17 
                 ~9-10 
                 ~7.9-9.6 
                 ~7.7-8.7 
               
               
                 Coefficient (10 −6   
               
               
                 m/m ° C.) 
               
               
                 Thermal Stress (MPa) 
                 552 
                 227 
                 355 
                 377 
                 314 
               
               
                 Thermal Stress/Y.S. 
                 ≦2.90 
                 ≦1.60 
                 ≦0.61 
                 ≦0.94 
                 ≦0.53 
               
               
                 Normalized cost to 
                 &gt;2 
                 &gt;2 
                 &gt;2 
                 1 
                 ~1 
               
               
                 HMS 
               
               
                   
               
            
           
         
       
     
     The low elastic modulus of Al alloys maintains the thermal stress at a fairly low level but still above the yield strength of the alloy. The high elastic modulus and thermal expansion coefficient of the  304  austenitic stainless steel and commercial cryogenic high Mn steel result in high thermal stress well above the yield strength of the materials. While the thermal stress of 9% Ni and the inventive high Mn steels (i.e., high Mn steels of the present disclosure) are comparable, the high strength of the inventive high Mn austenitic steels (i.e., high Mn steels of the present disclosure) makes it more capable as a material for LNG containment system and piping. Furthermore, the high Mn steels can be welded with strength matching welding consumables, a distinguished advantage over 9% Ni steel design which are derated due to the undermatched welds. Thus, the inventive high Mn austenitic steels of the present disclosure provide numerous advantages over the steels in the art. 
     In one aspect, the disclosure provides methods for improving the yield strength/cryogenic toughness/thermal stress of the steels through the control of microstructure and/or chemistry. In certain embodiments, the yield strength/cryogenic toughness/thermal stress of the steel compositions of the present disclosure can be improved/increased through the control of microstructure and/or chemistry. Some such possible routes include promoting phase transformations (e.g., to martensite/epsilon phases), twinning during deformation, and/or introducing hard erosion resistant second phase particles to the compositions. 
     In another aspect, the present disclosure provides high manganese steels tailored to resist wear and/or deformation (e.g., having improved wear/deformation resistance properties) at cryogenic temperatures. In general, due to the high strength at cryogenic temperatures, the high manganese steels of the present disclosure have advantages/potential in applications where wear and/or deformation resistances at cryogenic temperatures are desired/required (e.g., LNG production, transportation and petrochemical applications). 
     In another aspect, the present disclosure provides high manganese steels tailored to resist wear and/or deformation (e.g., having improved wear/deformation resistance properties) when used in applications that require wear and/or deformation resistance during large temperature variations (e.g., from about −170° C. to about room temperature), such as the temperature range experienced by a LNG storage tank and/or LNG piping system when in use (e.g., about −170° C.) compared to when not in use (e.g., room temperature). These properties are addressed by the high manganese steels of the present application. 
     Any of the steel compositions as described or embraced by the present disclosure may be advantageously utilized in many systems/applications (e.g., piping systems, material conveying systems, fluids/solids transport systems, in mining operations, and/or as material for liquefied natural gas (LNG) storage systems, and piping for LNG onloading/offloading piping systems), particularly where wear and/or deformation resistances are important/desired at cryogenic temperatures. In exemplary embodiments, the systems/methods of the present disclosure provide for low-cost and high strength steels (e.g., to be utilized in various cryogenic LNG-related applications). 
     In another aspect, any of the piping systems, material conveying systems, fluids/solids transport systems, in mining operations, and/or as material for liquefied natural gas (LNG) storage systems, and/or piping for LNG onloading/offloading piping systems presented herein may comprise waffle-free or wrinkle-free LNG storage systems and expansion-loop-free piping for LNG onloading/offloading piping systems. 
     In another aspect, the high Mn steels from any of the embodiments presented herein exhibit a low coefficient of thermal expansion (CTE), low elastic modulus, and/or low bulk modulus. 
     Thermal stress due to thermal expansion and contraction of materials from room temperature (25° C.) to service temperature (−163° C.) (i.e., liquefied natural gas (LNG) storage and/or LNG onloading/offloading), σ thermal , can be calculated from the following equation: 
       σ thermal=   E·α·ΔT  
 
     where E is elastic modulus or Young&#39;s modulus, α is thermal expansion coefficient, and ΔT is the temperature range from room temperature to service temperature. 
     Cryogenic materials with lower yield strength, for which σ thermal &gt;σy, have to be designed with “waffle” or wrinkled design for membrane type LNG storage tank and with expansion loops for cryogenic piping in order to keep the thermal stress below the yield stress of the material. Another aspect of the present disclosure is that the inventive steels can eliminate the necessity of wrinkled design or expansion loops in cryogenic structures. 
     In another embodiment, the inventive high Mn austenitic steel provides high yield strength, σ y , which is higher than thermal stress for LNG applications. In other words, σ thermal &lt;σ y . Preferably, σ thermal &lt;x·σ y  where x is smaller than 1. For example, x can be 0.9, 0.85, 0.8, 0.75, 0.7, 0.65, 0.6, 0.55, or 0.5. Preferably, x is smaller than 0.8. More preferably, x is smaller than 0.7. Even more preferably, x is smaller than 0.6. 
     In another embodiment, the inventive high Mn austenitic steel provides low CTE, approximately lower than 12 (10 −6  m/m° C.), and low elastic modulus (e.g., elastic modulus lower than 210 GPa), whilst maintaining high strength (e.g., strength higher than 590 MPa). 
     In another embodiment, the inventive high Mn austenitic steel comprises about 0.3% to about 1.2% Carbon, and from about 8% to 30% Manganese. 
     In yet another embodiment of the current disclosure, the inventive steel is alloyed with microalloying elements such as Nb, V, Ti, Mo, W and other alloying elements such as Cr, C, and N to enhance yield strength. In another embodiment of the present disclosure, the inventive steels can be welded with Ni-base weld wire (e.g., 625 Ni alloy) and/or high Mn alloyed weld wire. 
     In yet another embodiment, the inventive high Mn austenitic steels have a predominantly austentic structure stabilized at cryogenic temperature. The typical M d30  temperature, the highest temperature at which a designated amount of martensite forms under 30% deformation conditions, of the inventive steels are below service temperature of LNG containment system. 
     As discussed in further detail below, the fabrication methods/systems of the present disclosure can include one or more of the following steps: (i) providing a high work hardening rate matrix, through transformation induced plasticity (“TRIP”) and/or twin-induced plasticity (“TWIP”); (ii) providing meta-stability to induce phase transformation during service; (iii) providing optimum hardness of martensite (e.g., to be controlled by dissolved carbon content, to provide required erosion resistance); (iv) the dispersion of second phase particles (e.g., carbides, quasi-crystals, etc.) of varying size ranges within the compositions; (v) utilization of advantageous thermo-mechanical controlled process (“TMCP”) fabrication steps/schemes (e.g., to achieve at least some of the steps above); and/or (vi) exemplary joining methods, such as solid state joining (e.g., Friction Stir Welding). 
     In general, the high manganese steels of the present disclosure are relatively inexpensive alloys, and have potential applications where wear and/or deformation at cryogenic temperatures or the like of working components is important. In certain embodiments, the steel compositions have from about 0.30 to about 0.70 weight % carbon, and from about 11 to about 20 weight % manganese. 
     In exemplary embodiments, the steel has a fully austenitic structure obtained by quenching from a temperature above about 1000° C. In this condition, the hardness of the material is relatively low. One particularly advantageous feature of the high manganese steel is the strong work hardening capability. Under impact or other mechanical stress, the surface layer can increase its hardness rapidly by martensitic transformation or twinning, whereas other portions/parts of the steel remain substantially soft and/or ductile. This combination of low cost and high work hardening rate makes these steels advantageously suitable to be applied as wear resistant piping material at cryogenic temperatures or the like. 
     In general, the present disclosure provides for steels that exhibit a combination of high strength and erosion resistance. Also, as the result of their good formability, the high manganese steels as described herein can be used in a variety of settings, including, mining and automotive applications. 
     As noted, the present disclosure relates to high manganese steel chemistry and/or microstructures tailored to achieve improved strength, toughness, elastic modulus, thermal expansion coefficient and/or thermal conductivity. In exemplary embodiments, surface grain refinement may take place in a surface layer of certain high Mn steels either prior to and/or during service/use (e.g., formed in-situ). For example, the grain refinement at the surface can result in the formation of a layer which possesses the unique combinations of high strength and hardness, high ductility, and/or high toughness. Such fine grained (e.g., about 100 nm layer in height) or ultrafine grained (e.g., about 10 nm layer in height) surface layer may be formed either prior to and/or during service/use (e.g., formed in-situ), and can impart the desired strength, toughness, elastic modulus, thermal expansion coefficient and/or thermal conductivity to the steel. 
     In exemplary embodiments, such fine grained (e.g., about 100 nm layer) or ultrafine grained (e.g., about 10 nm layer) surface layer may be formed prior to use/installation of the exemplary steel by such surface deformation techniques such as, without limitation, shot peening, laser shock peening, and/or surface burnishing. 
     Current practice provides that the mechanical loads against the pipes in some piping systems are not strong enough to cause the maximum work hardening of the steel. In exemplary embodiments, the present disclosure provides high manganese steel with improved wear resistance, which can provide pipes for piping systems with advantageous wear life expectancies. 
     Current practice also provides that the steel in material conveying systems (e.g., piping systems, heavy equipment, etc.) often wears and/or fails prematurely, which leads to significant repair, replacement and/or maintenance/production costs. In exemplary embodiments, the present disclosure provides high manganese steel with improved wear and/or deformation resistance at cryogenic temperatures, which can provide pipes for piping systems with advantageous life expectancies. 
     In additional exemplary embodiments, the present disclosure provides for ferrous based components/compositions containing manganese. In certain embodiments, the components/compositions include from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron. The components/compositions can also include one or more alloying elements, such as, without limitation, chromium, nickel, cobalt, molybdenum, niobium, copper, titanium, vanadium, nitrogen, boron and combinations thereof. Exemplary ferrous based components/compositions containing manganese (and optionally other alloying elements) are described and disclosed in U.S. Patent Pub. No. 2012/0160363, the entire contents of which is hereby incorporated by reference in its entirety. 
     Component Composition 
     In exemplary embodiments and as noted above, the ferrous based compositions include from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron. 
     As such, the manganese level in the compositions may range from about 5 to 40 wt % of the total component/composition. The carbon level in the component/composition may range from 0.01 to 1.2 wt % of the total component/composition. In general, iron constitutes the substantial balance of the component/composition. 
     The components/compositions can also include one or more alloying elements, such as, without limitation, chromium, aluminum, nickel, cobalt, molybdenum, niobium, copper, titanium, vanadium, nitrogen, boron, zirconium, hafnium and combinations thereof. Weight percentages are based upon the weight of the total component/composition. 
     Chromium may be included in the component from about 0 to about 30 wt % (more preferably from 0.5 to 20 weight % of the total composition, even more preferably from 2 to 5 weight % of the total composition). Nickel may be included in the component from about 0 to about 20 wt % (more preferably from 0.5 to 20 weight % of the total composition, even more preferably from 1 to 5 weight % of the total composition). Cobalt may be included in the component from about 0 to about 20 wt % (more preferably from 0.5 to 20 weight % of the total composition, even more preferably from 1 to 5 weight % of the total composition). Aluminum may be included in the component from about 0 to about 15 wt % (more preferably from 0.5 to 10 weight % of the total composition, even more preferably from 1 to 5 weight % of the total composition. Molybdenum may be included in the component from about 0 to about 10 wt % (more preferably from 0.2 to 5 weight % of the total composition, even more preferably from 0.1 to 2 weight % of the total composition). Silicon may be included in the component from about 0 to about 10 wt % (more preferably from 0.1 to 6 weight % of the total composition, even more preferably from 0.1 to 0.5 weight % of the total composition). Niobium, copper, titanium and/or vanadium can each be included in the component from about 0.02 to about 10 wt % (more preferably from 0.02 to 5 weight % of the total composition, even more preferably from 0.02 to 2 weight % of the total composition). Nitrogen can be included in the component from about 0.01 to about 3.0 wt % (more preferably from 0.01 to 2.0 weight % of the total composition, even more preferably from 0.01 to 1.0 weight % of the total composition). Boron can be included in the component from about 0 to about 0.1 wt % (more preferably from 0.001 to 0.1 weight % of the total composition). 
     The ferrous based components/compositions containing manganese may also include another alloying element selected from the group consisting of zirconium, hafnium, and combinations thereof. Each of these other alloying elements may be included in the component/composition in ranges from about 0 to about 6 wt % (e.g., 0.2 to 5 wt %) based on the total weight of the component/composition. 
     In general, the mechanical properties of the high Mn steels of the present disclosure are dependent on the characteristics of strain-induced transformation, which is typically controlled by the chemical composition of the steels and/or the processing temperatures. Unlike conventional carbon steels, high Mn steels include a metastable austenite phase with a face centered cubic (fcc) structure at ambient temperature (e.g., 18-24° C.). 
     Upon straining, the metastable austenite phase can transform into several other phases through strain-induced transformation. More particularly, the austenite phase could transform into microtwins (fcc) structure (twin aligned with matrix), ε-martensite (hexagonal lattice), and α′-martensite (body centered tetragonal lattice), depending on steel chemistry and/or temperature. 
     These transformation products could impart a range of unique properties to high Mn steels. For example, fine microtwins effectively segment primary grains and act as strong obstacles for dislocation gliding. This leads to effective grain refinement which results in an excellent combination of high ultimate strength and ductility. 
     Chemical composition and temperature are known to be primary factors controlling the strain-induced phase transformation pathways as shown in  FIG. 1 . In general, high Mn steels can be divided into four groups depending on the stability of austenite phase upon straining and temperature, e.g., hilly stable (A), mildly metastable (B), moderately metastable (C) and highly metastable (D) Mn steel. The metastability of these phases is affected by both temperature and strain. These steels would tend to be more metastable (e.g., higher tendency to transform) at lower temperatures and higher strains. 
       FIG. 2  is an exemplary diagram of the phase stability and deformation mechanism of high Mn steels as a function of alloy chemistry and temperature. The letters (A, B, C, and D) indicate the various methods of transformation during deformation. In this diagram, steel A would deform by slip (similar to other metals and alloys), while steels B-D would transform during deformation. 
     Steel in area A, with high Mn content (e.g., greater than or equal to about 25 wt %), has stable austenite and deforms primarily by dislocation slip upon mechanical straining. In general, steels with a fully stabilized austenitic structure show lower mechanical strength but remain tough at cryogenic temperatures, provide low magnetic permeability and are highly resistant to hydrogen embrittlement. 
     Steel in area B, which is mildly metastable, can be produced with intermediate manganese content (e.g., from about 15 to about 25 wt % Mn, and about 0.6 wt % C). These steels form twins during deformation. A large amount of plastic elongation can be achieved by the formation of extensive deformation twins along with dislocation slip, a phenomenon known as Twinning-Induced Plasticity (TWIP). Twinning causes a high rate of work hardening as the microstructure is effectively refined, as the twin boundaries act like grain boundaries and strengthen the steel due to the dynamic Hall-Petch effect. TWIP steels combine extremely high tensile strength (e.g., greater than 150 ksi) with extremely high uniform elongation (e.g., greater than 95%), rendering them highly attractive for many applications. 
     The moderately metastable steels (Steel in area C) can transform into ε-martensite (hexagonal lattice) upon straining. Upon mechanical straining, these steels would deform predominantly by the formation of ε-martensite, along with dislocation slip and/or mechanical twinning. 
     The highly metastable steels (Steel in area D) will transform to a strong body-centered cubic phase (referred to as α′-martensite) upon deformation. This strong phase provides resistance to erosion resulting from the impingement of external, hard particles. Since the impact of the external particles results in the deformation of the near surface regions of the steel, these surface regions will transform during service, thereby providing resistance to erosion. Therefore, these steels have a “self-healing” characteristic in the sense that if the hard surface layer gets damaged, it would reform by the impact of the service. 
     Thus, the chemistry of the high Mn steels can be tailored to provide a range of properties (e.g., wear resistance, cryogenic toughness, high formability, erosion resistance) by controlling their transformation during deformation. 
     Other Alloying Concepts in High Mn Steels 
     Alloying elements in high Mn steels determine the stability of the austenite phase and strain-induced transformation pathways. In general, manganese is the main alloying element in high Mn steels, and it is important in stabilizing the austenitic structure both during cooling and deformation. In the Fe—Mn binary system, with increasing Mn content, the strain induced phase transformation pathway changes from α′-martensite to ε-martensite and then to micro-twinning. 
     Carbon is an effective austenite stabilizer and the carbon solubility is high in the austenite phase. Therefore, carbon alloying can be used to stabilize the austenite phase during cooling from the melt and during plastic deformation. Carbon also strengthens the matrix by solid solution hardening. As noted, the carbon in the components/compositions of the present disclosure may range from about 0.01 to about 1.2 wt % of the total component/composition. 
     Aluminum is a ferrite stabilizer and thus destabilizes austenite phase during cooling. The addition of aluminum to high Mn steels, however, stabilizes the austenite phase against strain-induced phase transformation during deformation. Furthermore, it strengthens the austenite by solid solution hardening. The addition of aluminum also enhances the corrosion resistance of the high Mn containing ferrous based components disclosed herein due to its high passivity. The aluminum in the components/compositions of the present disclosure may range from about 0.0 to about 15 wt % of the total component. 
     Silicon is a ferrite stabilizer and sustains the α′-martensite transformation while promoting ε-martensite formation upon deformation at ambient temperature. Due to solid solution strengthening, addition of Si strengthens the austenite phase by about 50 MPa per 1 wt % addition of Si. The silicon in the components/compositions of the present disclosure may range from about 0.01 to about 10 wt % of the total component. 
     Chromium additions to high Mn steel alloys enhance the formation of ferrite phase during cooling and increase corrosion resistance. Furthermore, the addition of Cr to the Fe—Mn alloy system reduces the thermal expansion coefficient. The chromium in the components/compositions of the present disclosure may range from about 0.5 to about 30 wt %. of the total component. 
     Based on the understanding of these alloying element effects on strain-induced phase transformation, suitable steel chemistries can be designed for specific applications. Some criteria for the design of high Mn steels can be the critical martensite transformation temperatures, e.g., M s  and M εs . M s  is a critical temperature below which austenite to α′-martensite transformation occurs, and M εs  is a critical temperature below which austenite to ε-martensite transformation takes place. 
     The effects of alloying elements on M s  and M εs  can be expressed as follows (unit of alloying elements in weight percent, and where A 3  is a critical temperature above which all ferrite phases (including α′- and ε-martensite phases) transform to austenite): 
       M s (K)=A 3 -410-200(C+1.4N)-18Ni-22Mn-7Cr-45Si-56Mo; and 
       M εs (K)=670-710(C+1.4N)-19Ni-12Mn-8Cr+13Si-2Mo-23Al 
     In general, only austenite to α′-martensite transformation takes place if M s  is much higher than M εs . If M εs  is much higher than M s , only austenite to ε-martensite transformation takes place. Both α′-martensite and ε-martensite phase transformation occur if M s  and M εs  are close to each other. 
     It is noted that the ferrous based components/compositions containing manganese may be utilized in a wide variety of applications/uses/systems (e.g., piping systems, material conveying systems, fluids/solids transport systems, in mining operations, and/or as material for liquefied natural gas (LNG) storage systems, and piping for LNG onloading/offloading piping systems, materials for cryogenic gas (e.g., Liquefied petroleum gas, Liquefied Ethylene gas, supercritical CO2) storage and piping, and cold liquid processing vessels and piping including amine scrubber, and high pressure heat exchangers. 
     For example, as described and disclosed in U.S. Patent Pub. No. 2012/0160363 noted above, the ferrous based components/compositions containing manganese of the present disclosure may find numerous non-limiting uses/applications in the oil, gas and/or petrochemical industry or the like (e.g., cryogenic applications, corrosion resistant applications, erosion resistant applications, natural gas liquefaction/transportation/storage type structures/components, oil/gas well completion and production structures/components, subterraneous drilling equipment, oil/gas refinery and chemical plant structures/components, coal mining structures/equipment, coal gasification structures/equipment, etc.). 
     The relatively low alloying content (e.g., less than about 20 wt % Mn, and about 0.6 wt % C) produces the highly metastable austenite phase. The highly metastable austenite phase often transforms into hard α′-martensite upon straining, which typically is an irreversible transformation. Upon surface wear of these steels, a surface layer of the highly metastable austenite phase can transform to α′-martensite phase. This friction-induced phase transformation leads to the formation of a thin, hard surface layer composed of martensite over an interior that consists of tough, untransformed austenite. This unique combination renders high Mn steels suitable for wear/erosion and impact resistant applications. 
     Moreover, the joining of the high Mn steels of the present disclosure can be performed using conventional (e.g., fusion, resistance welding, etc.) and emerging joining methods (e.g., laser, electron beam, friction stir welding, etc.), as described and disclosed in U.S. Patent Pub. No. 2012/0160363 noted above. In exemplary embodiments, preferred joining methods include solid state welding methods (e.g., resistance welding, friction stir welding), where such welding methods do not require the use of a weld metal, although the present disclosure is not limited thereto. 
     Bulk Modification 
     In exemplary embodiments, bulk modification is utilized to promote phase transformation (TRIP) and twinning (TWIP) during deformation. In general, the dispersed particles strengthen the materials/compositions, but have complex effects. It is noted that the dispersed particles may influence the: (i) chemistry of the composition matrix itself, (ii) grain size, and (iii) overall material/composition toughness. In general, the proper balance of these effects is important to exemplary embodiments of the present disclosure. 
     High manganese steel generally has a rapid work hardening rate because of the TRIP/TWIP effects. Their activations are typically triggered by the value of the stacking fault energy (“SFE”) of the alloy. It is noted that the plastic deformation is associated with martensitic transformation at low SFE values (e.g., less than about 12 mJ/m 2 ), and by twinning at intermediate SFE values. At even higher SFE values (e.g., greater than 35mJ/m 2 ), plasticity and strain hardening is typically controlled solely by dislocation sliding. As such, the SFE value is an important parameter in steel design. 
     The SFE is a function of alloy chemistry and temperature. The intrinsic stacking fault can be represented as a ε-martensite embryo of two planes in thickness. The SFE includes both volume energy and surface energy contributions. It has been demonstrated that the chemistry and temperature dependence of SFE arises largely from the volume energy difference between ε-martensite and austenite. Moreover, the volume free energy of phases can be obtained from available databases or the like. 
       FIG. 3  shows the predicted SFE values when adding each alloying element to FeMn13C0.6. Stated another way,  FIG. 3  displays the predicted influence of alloying elements on the SFE values of the FeMn13C0.6 reference and the deformation mechanism. 
     As shown in  FIG. 3 , the SFE contribution from the addition of various alloying elements is different. Carbon has the strongest effect, and manganese has the smallest influences. When the interaction of multiple alloying elements is considered, the dependence of SFE on chemistry will be complex and non-monotonic. In general, the deformation mechanism can be controlled by properly tailoring the bulk chemistry. 
     Second Phase Particle Dispersion Strengthening 
     In exemplary embodiments, the systems/methods of the present disclosure also include the introduction of second phase particles to further improve the wear resistance of the exemplary steel compositions. In certain non-limiting embodiments the exemplary systems/methods are described primarily with respect to carbide/nitride particles. However, it is noted that the systems/methods of the present disclosure may utilize, apply to and/or include other particles/precipitates, such as, without limitation, borides and oxides. In exemplary embodiments, when primarily carbide/nitride and oxide particles are considered, the grain size refinement can be an additional benefit from the second phase particles. 
     In general, the size and spatial distribution of the particles are important. It has been demonstrated that the effectiveness of the particles on the steel/material strengthening increases with decreasing particle size. Thus, fine particles generally contribute to the material wear resistance largely by strengthening the materials, while coarse particles typically provide additional resistance to erosive damage. 
     It is noted that the size and/or spatial distribution of the particles can be adjusted or optimized based on materials service conditions. For example, for compositions for use in a piping system or the like, the wear damage may be caused by sands having wide particle size distribution. Therefore, a bimodal particle distribution could be considered for the steel composition. It is noted that the fabrication or manufacture of high manganese steel with various type and size second phase particle can be achieved through various exemplary thermo-mechanical controlled processes (“TMCP”), as discussed further below. 
     In certain embodiments, the carbide/nitride precipitation can also locally enhance the TRIP or TWIP effects in the austenitic matrix. The interstitial elements (carbon and nitrogen) concentration in carbide/nitride particles is much higher than the average value of the steel. Due to diffusion gradients at the interface, the interstitial elements could be depleted in precipitates surrounding the matrix, which results in a lower activation energy for TRIP or TWIP. 
       FIG. 4  (schematic, not in scale) displays the overall effect of carbide precipitates on mechanical properties and the corresponding deformation mechanism. Compared to the fully austenitic steel with substantially the same chemistry, the high manganese steel with the exemplary carbide/nitride particles can have a higher yield strength and work hardening capability. In exemplary embodiments, the combination of hard particles and a work-hardenable material matrix makes the compositions of the present disclosure suitable to withstand and/or reduce the abrasive wear effects caused by operational use (e.g., by hard particle cutting/shearing or the like). 
     Fabrication and Microalloying 
     The steel compositions/components of the present disclosure can be fabricated or manufactured by various processing techniques including, but not limited to, various exemplary thermo-mechanical controlled processing (“TMCP”) techniques, steps or methods. In general, some TMCP processes have been utilized to produce low alloy steel, particularly where grain size and microstructure refinement is desired. 
     In exemplary embodiments, to produce the desired carbide/nitride particles, the particles should be in a substantially dissolved state before the deformation, as undissolved particles will suffer relatively rapid coarsening at the elevated temperatures. The controlled deformation should take place below the recrystallization stop temperature so that deformation results in elongated austenite grains filled with intra-granular crystalline defects, which are the preferred sites for nucleation. 
     A slow cooling or isothermal holding is then required to promote the particles precipitation. Finally, a rapid quench is applied to keep a fully austenitic matrix. 
       FIG. 5  illustrates an exemplary fabrication schedule for the production of steel compositions/components according to the present disclosure. As such,  FIG. 5  is a schematic drawing of an exemplary steel fabrication method of the present disclosure. As shown in  FIG. 5 , T nr  is the austenite recrystallization stop temperature, and A s  is the austenite start temperature. 
     In exemplary embodiments, the TMCP methods have a synergistic effect of micro-alloy additions. Depending on the alloying elements to be added/utilized in the composition, the appropriate thermo-mechanical conditions should be selected in order to produce the desired fine particles. In general, the alloying elements utilized in the methods of the present disclosure can have some effect on either the TMCP, or the bulk property modification, or both. 
     In certain embodiments, carbon is the one of the most effective alloying elements to control the bulk deformation mechanism, promote carbide precipitation and stabilize the austenite phase during cooling. It is noted that the total carbon content of the compositions could be much larger or higher compared to conventional high manganese steel, but the amount of carbon in solution after TMCP steps should be controlled to a much lower level. 
     In exemplary embodiments, manganese is the austenite phase stabilizer. This element can be mainly added to the compositions to maintain a fully austenitic matrix during cooling and TMCP. In general, it has little effect on the deformation mechanism. 
     Chromium is a carbide former. It will promote different types of carbide, such as M7C and M23C6, depending on the alloy level and/or thermal treatment temperature. Moreover, chromium addition is typically important for corrosion resistance enhancement. 
     Niobium and titanium are effective elements to retard the recrystallization during TMCP by forming strain induced (e.g., Ti, Nb C, N) precipitation on the deformed austenite. In addition, the niobium and/or titanium addition facilitates the bulk carbon concentration modification according to exemplary embodiments of the present disclosure. 
     Aluminum and silicon are added to tune or adjust the SFE of the high manganese steel of the present disclosure. It is noted that aluminum addition can facilitate quasi-crystalline phase formation, as discussed below. 
     Quasi-Crystal Precipitation Hardened High Mn Steels 
     It is another object of the present disclosure to provide high Mn steels utilizing precipitation hardening of quasi-crystals. In exemplary embodiments, high Mn steels can be strengthened by the precipitation of quasi-crystals, and such structures can be achieved by heat treating at elevated temperatures (e.g., up to about 700° C.). 
     In general, quasi-crystalline materials have periodic atomic structures (e.g., 5-fold or 10-fold rotational symmetry), but usually do not conform to the 3-D symmetry typical of ordinary crystalline materials. Due to their crystallographic structure, quasi-crystalline materials with tailored chemistry exhibit unique properties, which are attractive for the strengthening of high Mn steels. 
     It is noted that the quasi-crystalline precipitates can provide higher strengthening effects than that of crystalline precipitates (e.g., carbides), because of the difficulty of dislocations to move through quasi-crystal lattices. Furthermore, quasi-crystals usually will not grow beyond certain sizes unlike crystalline precipitates, thereby alleviating over-aging concerns associated with certain crystalline precipitates. 
     Quasi-crystal materials typically provide non-stick surface properties due to their low surface energy (e.g., about 30mJ/m 2 ) on stainless steel substrates in icosahedral Al—Cu—Fe chemistries. Due to their low surface energy, quasi-crystal materials exhibit a low friction coefficient (e.g., about 0.05) in scratch tests with diamond indentor in dry air, combined with relatively high micro-hardness. Quasi-crystalline materials are found in Al-TM (TM=transition metals; e.g., V, Cr, Mn), Al—(Mn, Cu, Fe)—(Si), and Al—Cu-TM (e.g., Cr, Fe, Mn, Mo) systems. 
     Ultrafine Grained High Mn Steels 
     In exemplary embodiments, improved steel compositions (e.g., ultrafine grained high Mn steel compositions) can be fabricated by exemplary thermo-mechanical controlled processes (TMCP). In certain embodiments, especially in lower Mn alloying chemistry such as 8 wt. % or less in which ferrite or martensite phase is thermodynamically more stable than the austenite phase, the TMCP of the present disclosure includes heavy plastic deformation at ambient (e.g., 18-24° C.) or cryogenic (e.g., −196° C.) or intermediate temperature to induce martensite formation, and subsequent annealing at elevated temperatures to reverse deformation-induced martensite into ultrafine grained austenite. An exemplary thermo-mechanical controlled process is schematically shown in  FIG. 6 .  FIG. 6  depicts a schematic of an exemplary fabrication method for producing ultrafine grained high Mn steels according to the present disclosure. As shown in  FIG. 6 , Af is the austenite finish temperature (austenite recrystallization stop temperature), and A s  is the austenite start temperature. 
     In exemplary embodiments and after heating and holding the steel composition at a normalizing temperature, the metastable austenite phase of the steel composition is transformed to a strain-induced martensite phase by heavy plastic deformation at ambient (e.g., 18-24° C.) or cryogenic (e.g., −196° C.) or intermediate temperatures ( FIG. 6 ). The strain-induced martensite phase may be further heavily deformed to destroy lath or plate structures prior to a reversion treatment (e.g., reversion annealing in  FIG. 6 ). The strain-induced martensite phase may be reverted to the austenite phase at temperatures low enough to suppress the grain coarsening of the reverted austenite phase. In exemplary embodiments, the chemistry of the steel compositions of the present disclosure (e.g., high Mn steel compositions) can be tailored so that the martensite start temperature (M s ) of reverted austenite is below room temperature (e.g., 18-24° C.). 
     Exemplary Methods for Fabrication 
     The present disclosure provides for a method for fabricating a ferrous based component including: a) providing a composition having from about 5 to about 40 weight % manganese, preferably from about 9 to about 25 weight % manganese, even more preferably from about 12 to about 20 weight % manganese, and from about 0.01 to about 1.2 weight % carbon, preferably from about 0.3% to about 1.2 weight % carbon, even more preferably from about 0.3% to about 0.7 weight % carbon, and the balance iron, b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition (e.g., to a temperature to homogenize the composition); c) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and d) quenching or accelerated cooling the composition. 
     The present disclosure provides for a method for fabricating a ferrous based component wherein after step d), the matrix of the composition is predominantly or substantially in the austenitic phase. In one or more embodiments, the volume percent of austenite in the steel composition is from about 50 wt % to about 100 wt %, more preferably from about 80 wt % to about 99 wt %. 
     The steel composition is preferably processed into predominantly or substantially austenitic plates using a hot rolling process. In one or more embodiments, a steel billet/slab from the compositions described is first formed, such as, for example, through a continuous casting process. The billet/slab can then be re-heated to a temperature (“reheat temperature”) within the range of about 1,000° C. to about 1,300° C., more preferably within the range of about 1050° C. to 1250° C., even more preferably within the range of about 1100° C. to 1200° C. Preferably, the reheat temperature is sufficient to: (i) substantially homogenize the steel slab/composition, (ii) dissolve substantially all the carbide and/or carbonitrides, when present, in the steel slab/composition, and (iii) establish fine initial austenite grains in the steel slab/composition. 
     The re-heated slab/composition can then be hot rolled in one or more passes. In exemplary embodiments, the reheated slabs/billets can be cooled to the rolling start temperature. In one or more embodiments, the rolling start temperature is above 1100° C., preferably above 1080° C., even more preferably above 1050° C. 
     In exemplary embodiments, the final rolling for plate thickness reduction can be completed at a “finish rolling temperature”. In one or more embodiments, the finish rolling temperature is above about 700° C., preferably above about 800° C., more preferably about 850° C. Thereafter, the hot rolled plate can be cooled (e.g., in air) to a first cooling temperature or accelerated cooling start temperature (“ACST”), at which an accelerated cooling starts to cool the plates at a rate of at least about 10° C. per second to a second cooling temperature or accelerated cooling finish temperature (“ACFT”). After the cooling to the ACFT, the steel plate/composition can be cooled to room temperature (e.g., ambient temperature) in ambient air. Preferably, the steel plate/composition is allowed to cool on its own to room temperature. 
     In one or more embodiments, the ACST is about 750° C. or more, about 800° C. or more, about 850° C. or more, or about 900° C. or more. In one or more embodiments, the ACST can range from about 700° C. to about 1000° C. In one or more embodiments, the ACST can range from about 750° C. to about 950° C. Preferably, the ACST ranges from a low of about 650° C., 700° C., or 750° C. to a high of about 900° C., 950° C., or 1000° C. In one or more embodiments, the ACST can be about 750° C., about 800° C., about 850° C., about 890° C., about 900° C., about 930° C., about 950° C., about 960° C., about 970° C., about 980° C., or about 990° C. 
     In one or more embodiments, the ACFT can range from about 0° C. to about 500° C. Preferably, the ACFT ranges from a low of about 0° C., 10° C., or 20° C. to a high of about 150° C., 200° C., or 300° C. 
     Without being bound by any theory, it is believed that the rapid cooling (e.g., more than about 10° C./sec cooling rate) to the low accelerated cooling finish temperature (“ACFT”) retards at least a portion of the carbon and/or nitrogen atoms from diffusing from the austenite phase of the steel composition to the grain boundary or second phase. It is further believed that the high accelerated cooling start temperature (“ACST”) retards at least a portion of the carbon and/or nitrogen atoms from forming precipitates such as, for example, carbides, carbonitrides, and/or nitrides during subsequent cooling to the ACFT. As such, the amount of precipitates at the grain boundaries is reduced. Therefore, the steel&#39;s fracture toughness and/or resistance to cracking is enhanced. 
     Following the rolling and cooling steps, the plate can be formed into pipes or the like (e.g., linepipe). Any suitable method for forming pipe can be used. Preferably, the precursor steel plate is fabricated into linepipe by a conventional UOE process or JCOE process which is known in the art. 
     EXAMPLES 
     The present disclosure will be further described with respect to the following examples; however, the scope of the disclosure is not limited thereby. The following examples illustrate improved systems and methods for fabricating or producing improved steel compositions (e.g., improved high Mn steel compositions having enhanced wear and/or deformation resistance at cryogenic temperatures or the like). As illustrated in the below examples, the present disclosure illustrates that the advantageous steel compositions/components of the present disclosure improve one or more of the following properties: strength, toughness, elastic modulus, thermal expansion coefficient and/or thermal conductivity. In general, the yield strength/cryogenic toughness/thermal stress of the steels of the present disclosure can be improved/increased through the control of microstructure and/or chemistry. As noted, some possible routes include promoting phase transformations (e.g., to martensite or epsilon phases), twinning during deformation, and/or introducing hard erosion resistant second phase particles to the compositions. 
     In exemplary embodiments, the present disclosure provides for a ferrous based component fabricated according to the steps comprising: a) providing a composition having from about 5 to about 40 weight % manganese, from about 0.01 to about 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     Example 1: Fabrication of High Mn Steels of the Present Disclosure 
     In one exemplary embodiment of the disclosure, steel plates having the chemistry shown in Table 2 are fabricated by vacuum induction melting and hot rolling. Steel plates with 25 mm thickness were fabricated by finish rolling at around 850° C. followed by accelerated cooling to room temperature. 
     
       
         
           
               
             
               
                 TABLE 2 
               
             
            
               
                   
               
               
                 Chemistry of inventive steels (weight percent) 
               
            
           
           
               
               
               
               
               
               
               
               
               
               
               
            
               
                 Steel 
                 C 
                 Si 
                 Mn 
                 Cr 
                 Mo 
                 Ti 
                 Nb 
                 Cu 
                 V 
                 N 
               
               
                   
               
               
                 EM501 
                 0.605 
                 0.150 
                 17.53 
                 3.12 
                 — 
                 — 
                 — 
                 0.50 
                 — 
                 — 
               
               
                 EM502 
                 0.600 
                 0.132 
                 18.30 
                 3.03 
                 — 
                 — 
                 0.024 
                 0.50 
                 — 
                 — 
               
               
                 EM503 
                 0.611 
                 0.138 
                 18.03 
                 3.01 
                 — 
                 0.021 
                 0.023 
                 0.50 
                 — 
                 0.015 
               
               
                 EM504 
                 0.620 
                 0.154 
                 18.10 
                 3.01 
                 — 
                 0.017 
                 0.023 
                 0.50 
                 — 
                 0.022 
               
               
                 EM505 
                 0.600 
                 0.124 
                 18.20 
                 2.98 
                 — 
                 — 
                 0.020 
                 0.50 
                 0.097 
                 — 
               
               
                 EM506 
                 0.587 
                 0.160 
                 18.02 
                 2.98 
                 0.52 
                 0.083 
                 0.022 
                 0.50 
                 — 
                 — 
               
               
                   
               
            
           
         
       
     
     Example 2: Mechanical Properties of High Mn Steels of the Present Disclosure 
     Tables 3 and Table 4 show the mechanical properties of inventive steels at cryogenic temperature and ambient temperature (around 25° C.), respectively. All of the steel plates were characterized by minimum yield strength ranging from about 600 MPa to 630 MPa at ambient temperature in the as-rolled condition without cold deformation. 
     
       
         
           
               
             
               
                 TABLE 3 
               
             
            
               
                   
               
               
                 Cryogenic mechanical properties of inventive steels. 
               
            
           
           
               
               
               
            
               
                   
                 Cryogenic Mechanical Properties @ −164° C. 
                 Charpy 
               
            
           
           
               
               
               
               
               
               
               
            
               
                   
                 0.2% Yield 
                 Ultimate 
                 Elastic 
                   
                 Reduction 
                 Impact 
               
               
                 Sample 
                 Strength 
                 Tensile Strength 
                 Modulus 
                 Elongation 
                 of Area 
                 Energy @ 
               
               
                 ID 
                 (MPa) 
                 (MPa) 
                 (GPa) 
                 (%) 
                 (%) 
                 −196° C. (J) 
               
               
                   
               
               
                 EM501 
                 880 
                 1508 
                 185.5 
                 48 
                 37 
                 67.3 
               
               
                 EM502 
                 802 
                 1500 
                 177.9 
                 48 
                 34 
                 76.8 
               
               
                 EM503 
                 821 
                 1489 
                 180.0 
                 50 
                 48 
                 73.2 
               
               
                 EM504 
                 799 
                 1518 
                 184.8 
                 49 
                 46 
                 74.1 
               
               
                 EM505 
                 776 
                 1499 
                 181.3 
                 50 
                 48 
                 70.5 
               
               
                 EM506 
                 732 
                 1417 
                 160.0 
                 45 
                 30 
                 56.9 
               
               
                   
               
            
           
         
       
     
     
       
         
           
               
             
               
                 TABLE 4 
               
             
            
               
                   
               
               
                 Ambient temperature mechanical properties of inventive steels. 
               
            
           
           
               
               
               
            
               
                   
                 Room Temperature Mechanical Properties @ 25° C. 
                 Charpy 
               
            
           
           
               
               
               
               
               
               
               
            
               
                   
                 0.2% Yield 
                 Ultimate 
                 Elastic 
                   
                 Reduction 
                 Impact 
               
               
                 Sample 
                 Strength 
                 Tensile Strength 
                 Modulus 
                 Elongation 
                 of Area 
                 Energy @ 
               
               
                 ID 
                 (MPa) 
                 (MPa) 
                 (GPa) 
                 (%) 
                 (%) 
                 −40° C. (J) 
               
               
                   
               
            
           
           
               
               
               
               
               
               
               
            
               
                 EM501 
                 598 
                 1085 
                 143.0 
                 57 
                 — 
                 119.4 
               
               
                 EM502 
                 611 
                 1077 
                 150.2 
                 55 
                 — 
                 112.0 
               
               
                 EM503 
                 618 
                 1086 
                 156.5 
                 61 
                 — 
                 112.8 
               
               
                 EM504 
                 630 
                 1090 
                 134.8 
                 55 
                 — 
                 128.7 
               
               
                 EM505 
                 625 
                 1081 
                 147.6 
                 54 
                 — 
                 106.9 
               
               
                 EM506 
                 624 
                 1095 
                 157.6 
                 54 
                 — 
                 89.8 
               
               
                   
               
            
           
         
       
     
     Table 5 shows the average coefficient of thermal expansion (CTE), and mechanical properties of inventive steels. The measurement of CTE were made with a LVDT (linear variable differential transducer)-based quartz dilatometer measurement system (ASTM standard E228). The change of specimen length has been recorded as the temperature was cycled between −196° C. and 25° C. at a maximum rate of 2° C./minute. T-type thermocouples were used for temperature measurement in the dilatometer. The CTE (α) were calculated as follows: 
     
       
         
           
             
               α 
               
                 t 
                  
                 
                     
                 
                  
                 12 
               
             
             = 
             
               
                 { 
                 
                   
                     ( 
                     
                       
                         L 
                         
                           t 
                            
                           
                               
                           
                            
                           2 
                         
                       
                       - 
                       
                         L 
                         
                           t 
                            
                           
                               
                           
                            
                           1 
                         
                       
                     
                     ) 
                   
                   / 
                   
                     L 
                     0 
                   
                 
                 } 
               
               
                 
                   T 
                   
                     t 
                      
                     
                         
                     
                      
                     2 
                   
                 
                 - 
                 
                   T 
                   
                     t 
                      
                     
                         
                     
                      
                     1 
                   
                 
               
             
           
         
       
     
     where L t1  and L t2  are the measured displacements of the specimens at time t 1 , and t 2  respectively; 
     T t1  and T t2  are the measured temperatures of the specimens at time t 1 , and t 2  respectively; and 
     L 0  is the initial specimen length 
     The average CTE values were calculated as a secant slope based of the end points of the polynomial regression from −196° C. and 25° C. Elastic modulus at cryogenic temperature and 0.2% yield stress at ambient temperature were used for the calculation of thermal stress using Equation (1). The inventive steels showed thermal stress lower than 50% of their yield strength. 
     
       
         
           
               
             
               
                 TABLE 5 
               
             
            
               
                   
               
               
                 Average coefficient of thermal expansion and 
               
               
                 mechanical properties of inventive steels. 
               
            
           
           
               
               
               
               
               
               
            
               
                   
                 Average 
                   
                   
                   
                   
               
               
                   
                 Thermal 
                   
                   
                 0.2% 
                 Thermal 
               
               
                   
                 Expansion 
                 Elastic 
                 Thermal 
                 Yield 
                 Stress/ 
               
               
                   
                 Coefficient 
                 Modulus 
                 Stress 
                 Strength 
                 0.2% Yield 
               
               
                 Steel 
                 (10 −6  m/m ° C.) 
                 (GPa) 
                 (MPa) 
                 (MPa) 
                 Strength 
               
               
                   
               
               
                 EM501 
                 8.5 
                 185.5 
                 299.5 
                 598 
                 0.50 
               
               
                 EM502 
                 8.1 
                 177.9 
                 273.8 
                 611 
                 0.45 
               
               
                 EM503 
                 8.5 
                 180.0 
                 290.6 
                 618 
                 0.47 
               
               
                 EM504 
                 8.3 
                 184.8 
                 291.4 
                 630 
                 0.46 
               
               
                 EM505 
                 7.9 
                 181.3 
                 272.2 
                 625 
                 0.44 
               
               
                 EM506 
                 8.1 
                 160.0 
                 246.2 
                 624 
                 0.46 
               
               
                   
               
            
           
         
       
     
     Without being bound by any theory, it is noted that an increase in carbon content in the high Mn steel matrix enhances the mechanical strength of the steel and the work hardening rate. In exemplary embodiments, Mn alloying in the range of about 12-20 weight % of the total composition stabilizes the austenite phase, and increases the carbon solubility in the steel matrix. The Cr alloying up to about 3 weight % increases the corrosion resistance and the mechanical strength by solution strengthening. The Cu alloying of about 0.5 to about 2 weight % increases carbon solubility and corrosion resistance. 
     In exemplary embodiments, the (TMCP) hot rolling parameters can be adjusted to obtain steel compositions having a refined grain sizes of about 200 μm or less, and/or low carbide precipitate fractions of about 5 volume % or less. The methods may include a finish rolling step or steps at lower temperatures, which would introduce deformation banding/dislocations tangles to thereby enhance the formation of fine intra-grain precipitates. 
     The exemplary modified TMCP hot rolling steps/parameters can be combined with the addition of various micro-alloying elements such as, without limitation, V, Nb, Ti, Mo and/or N. It has been found that the micro-alloying elements in high Mn steels can result in the formation of fine carbide/nitride/carbo-nitride precipitates finely dispersed in the steel matrix. The finely dispersed precipitates can retard grain coarsening during reheating and recrystallization during hot rolling, which thereby advantageously enhances the strength of the steel compositions/components of the present disclosure. 
     PCT/EP Clauses 
     1. A method for fabricating a ferrous based component comprising: a) providing a composition having from 5 to 40 weight % manganese, from 0.01 to 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition or to a temperature to homogenize the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     2. The method of clause 1, wherein after step e), the carbide precipitate fraction volume of the composition is 5 volume % or less of the composition. 
     3. The method of clauses 1 or 2, wherein after step e), the composition has a microstructure having a refined grain size of 100 μm or less. 
     4. The method of clause 3, wherein the microstructure having a refined grain size of 100 μm or less includes a surface layer of the composition; wherein the thickness of the surface layer is from 10 nm to 5000 nm; and wherein the surface layer is formed prior to or during use of the composition. 
     5. The method of clause 4, wherein the surface layer is formed via a surface deformation technique selected from the group consisting of shot peening, laser shock peening, surface burnishing and combinations thereof. 
     6. The method of any one of clauses 1-5, wherein prior to step e), the composition is slowly cooled or isothermally held; and wherein step e) includes rapidly quenching the composition. 
     7. The method of any one of clauses 1-6, wherein step d) includes deforming the composition while the composition is at a temperature below the austenite recrystallization temperature and above the martensite transformation start temperature. 
     8. The method of any one of clauses 1-7, wherein step d) includes deforming the composition to induce martensite formation of the composition; wherein the composition is deformed at a temperature of from 18° C. to 24° C. to induce martensite formation of the composition; and further comprising, after step d), heating the composition to a temperature above the austenite recrystallization stop temperature; wherein heating the composition to a temperature above the austenite recrystallization stop temperature after step d) reverses deformation-induced martensite of the composition into ultrafine grained austenite; and wherein the martensite start temperature of the ultrafine grained austenite is below 24° C. 
     9. The method of any one of clauses 1-8, further comprising, after step e), heating the composition to a temperature above the austenite recrystallization stop temperature, and then quenching the composition. 
     10. The method of any one of clauses 1-9, further comprising, prior to step c), deforming the composition while the composition is at a temperature above the austenite recrystallization stop temperature. 
     11. The method of any one of clauses 1-10, wherein step c) includes cooling the composition at a rate of from 2° C. per second to 60° C. per second. 
     12. The method of any one of clauses 1-11, wherein the composition further includes one or more alloying elements selected from the group consisting of chromium, aluminum, silicon, nickel, cobalt, molybdenum, niobium, copper, titanium, vanadium, nitrogen, boron, zirconium, hafnium and combinations thereof. 
     13. The method of clause 12, wherein the chromium ranges from 0.5 to 30 weight % of the total composition; wherein each of the nickel or cobalt ranges from 0.5 to 20 weight % of the total composition; wherein the aluminum ranges from 0.2 to 15 weight % of the total composition; wherein each of the molybdenum, niobium, copper, titanium or vanadium ranges from 0.01 to 10 weight % of the total composition; wherein the silicon ranges from 0.1 to 10 weight % of the total composition; wherein the nitrogen ranges from 0.001 to 3.0 weight % of the total composition; wherein the boron ranges from 0.001 to 0.1 weight % of the total composition; and wherein each of the zirconium or hafnium ranges from 0.2 to 6 weight % of the total composition. 
     14. The method of any one of clauses 1-13, wherein the composition includes from 8 to 20 weight % manganese, from 0.30 to 0.7 weight % carbon, from 0.5 to 3 weight % chromium, from 0.5 to 2.0 weight % copper, from 0.1 to 1 weight % silicon, and the balance iron. 
     15. A ferrous based component fabricated according to the steps comprising: a) providing a composition having from 5 to 40 weight % manganese, from 0.01 to 1.2 weight % carbon, and the balance iron; b) heating the composition to a temperature above the austenite recrystallization stop temperature of the composition; c) cooling the composition to a temperature below the austenite recrystallization stop temperature of the composition; d) deforming the composition while the composition is at a temperature below the austenite recrystallization stop temperature of the composition; and e) quenching the composition. 
     Whereas the disclosure has been described principally in connection with steel compositions for use in components for material conveying systems, fluids/solids transport systems, mining operations, oil sand piping systems, earth-moving equipment, drilling components, and/or oil/gas/petrochemical applications, such descriptions have been utilized only for purposes of disclosure and are not intended as limiting the disclosure. To the contrary, it is to be recognized that the disclosed steel compositions are capable of use in a wide variety of applications, systems, operations and/or industries. 
     Although the systems and methods of the present disclosure have been described with reference to exemplary embodiments thereof, the present disclosure is not limited to such exemplary embodiments and/or implementations. Rather, the systems and methods of the present disclosure are susceptible to many implementations and applications, as will be readily apparent to persons skilled in the art from the disclosure hereof. The present disclosure expressly encompasses such modifications, enhancements and/or variations of the disclosed embodiments. Since many changes could be made in the above construction and many widely different embodiments of this disclosure could be made without departing from the scope thereof, it is intended that all matter contained in the drawings and specification shall be interpreted as illustrative and not in a limiting sense. Additional modifications, changes, and substitutions are intended in the foregoing disclosure. Accordingly, it is appropriate that the appended claims be construed broadly and in a manner consistent with the scope of the disclosure.