JP2005519203A - Nanocarbide precipitation strengthened ultra high strength corrosion resistant structural steel - Google Patents

Abstract

Nanocarbide precipitation strengthened ultra high strength corrosion resistant structural steel has a combination of strength and corrosion resistance, which is about 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co ), 0 to 10% nickel (Ni), 6 to 12% chromium (Cr), less than 1% silicon (Si), less than 0.5% manganese (Mn), and less than 0.15% copper (Cu) approximately less than 3% molybdenum (Mo), less than 0.3% niobium (Nb), less than 0.8% vanadium (V), less than 0.2% tantalum (Ta), 3% An additive selected from the group comprising less than tungsten (W), and combinations thereof, approximately less than 0.2% titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, 0 Less than 15% zirconium (Zr), 0.005 Less than boron (B), and additional additives selected from the group comprising combinations thereof, approximately less than 0.02% sulfur (S), 0.012% phosphorus (P), 0.015% oxygen In combination with (O) and 0.015% nitrogen (N) impurities, the balance being substantially iron (Fe), incidental elements and other impurities. This alloy is strengthened by nanometer-scale MZC carbide in a fine lath martensite matrix, and the enhanced chemical distribution of Cr from the matrix to the surface provides an oxide passivation film suitable for corrosion resistance. This alloy with UTS greater than 280 ksi is useful for applications such as aircraft landing gear, machinery and tools used in adverse environments, and other applications where ultra high strength corrosion resistant structural steel alloys are desirable.

Description

(Cross-reference with related applications)
This international application is based on US patent application Ser. No. 10 / 071,688 filed on Feb. 8, 2002, which is based on the following provisional application, which is hereby incorporated by reference and has priority to them: U.S. Patent Application No. 60/267, it on the 9th February 2001 application, entitled "Nano-Precipitation Strengthened Ultra-High Strength Corrosion Resistant Structural Steels," and n. US Patent Application No. 60, filed Sep. 21, 2001, entitled “Corrosion Resistant Structural Steels”. No. 323,996.

In a main aspect, the present invention relates to a cobalt, nickel, chromium stainless martensitic steel alloy having ultra high strength and corrosion resistance, characterized by nanoscale sized carbide precipitation, particularly M ₂ C precipitation.

  Major structural elements in aerospace and other high-performance structures are made almost exclusively of ultra-high strength steel due to the heavy weight, size and possibly cost penalties associated with the use of other materials . However, ultra high strength steels having a tensile strength in the range of at least 240 ksi to 300 ksi are inferior in general corrosion resistance and are prone to hydrogen and environmental embrittlement.

  Therefore, cadmium plating of elements is typically used to provide general corrosion resistance in aerospace and other structural steel components, and hard chrome plating is mainly used when wear resistance is required. These coatings are disadvantageous in terms of cost, manufacturing, environment and reliability. Thus, the goal in the design or discovery of ultra high strength steel alloys is to eliminate the need for cadmium and chromium coatings without mechanical defects or reduced strength. One of the achievements targeted by the subject invention alloys has a tensile strength greater than about 240 ksi, does not require a cadmium coating, and exhibits wear resistance without chrome plating or other protective and wear resistant coatings. It is to replace non-stainless structural steel with stainless steel or corrosion resistant steel.

One of the ultra-high strength steels most widely used in aerospace structural applications is 300M steel. This alloy is essentially 4340 steel that has been modified to obtain a slightly higher Stage I tempering temperature, thereby allowing the combustion of brittle hydrogen introduced during processing. The Aerospace Material Specification AMS 6257A [SAE International, Warrendale, PA, 2001], incorporated herein, covers most of the use of 300M steel in aerospace applications. In this specification, the minimum tensile properties are maximum tensile strength (UTS) 280 ksi, yield strength (YS) 230 ksi, elongation 8% and area reduction of 30%. The average plane strain mode I fracture toughness incorporated here is

[Philip, T. et al. V. and T.A. J. et al. McCaffrey, Ultrahigh-Strength Steels, Properties and Selection: Irons, Steels, and High-Performance Alloys, Materials Park, OH, 90, 4: Internal 30. Stress corrosion cracking resistance in 3.5 wt% sodium chloride aqueous solution is

It is reported.

  The high tensile strength of 300M steel allows the design of lightweight structural components in aerospace systems such as landing gear. However, the general lack of corrosion resistance requires a cadmium coating and low stress corrosion cracking resistance results in significant outdoor failure due to environmental embrittlement.

  The precipitation-strengthened stainless steel, primarily 15-5PH, [AMS 5659K, SAE International, Warrendale, PA, 1998], incorporated herein, can also be used in structural aerospace elements, but typically its low strength. Therefore, it can be used only in low-load applications where the weight penalty is not large. Corrosion resistance is sufficient for such alloys and therefore can eliminate cadmium plating, but the minimum tensile properties of 15-5PH at maximum strength H900 conditions are only 190 ksi UTS and 170 ksi YS. This limits its application to components that are not strength limited.

  Another precipitation-strengthened stainless steel, Carpenter Custom 465 ™ steel [Alloy Digest, SS-716, Materials Park, OH, ASM International, 1998], incorporated here, uses intermetallic precipitation and is slightly below 270 ksi. The maximum UTS is reached. At that strength level, the Custom 465 ™ steel incorporated herein has a low Charpy V-notch impact energy of about 5 ft-lb [Kimmel, W. et al. M.M. , N.M. S. Kuhn et al., Cryogenic Model Materials, 39th AIAA Aerospace Sciences Meeting & Exhibit, Reno, NV, 2001]. For most structural applications, Custom 465 ™ steel must be used in conditions that limit its UTS to less than 270 ksi in order to maintain adequate Charpy V-notch impact resistance.

  Several secondary hardened stainless steels have been developed that reach the highest strength levels up to 270 ksi. These are described in US Pat. 26,225, 3,756,808, 3,873,378, and 5,358,577. These stainless steels use higher chromium levels to maintain corrosion resistance, thus compromising strength. The main feature of these alloys is the large amount of austenite that remains and forms during secondary hardening. This austenite alters the flow behavior of the alloys, and although they can achieve UTS as strong as 270 ksi, their yield strength is less than 200 ksi. This large disparity between yield and maximum limits the applications in which these steels can be used. Accordingly, there remains a need for an ultra high strength non-corrosive steel alloy having a yield strength of at least about 230 ksi and a maximum tensile strength of at least about 280 ksi.

(Summary of the Invention)
Briefly, the present invention is about 0.1 to 0.3% carbon (C), 8 to 17% cobalt (Co), less than 10% nickel (Ni), greater than 6% by weight, And other elemental additives including less than 11% chromium (Cr), and less than 3% molybdenum (Mo) in small amounts of Si, Cu, Mn, Nb, V, Ta, W, Ti, Zr, rare earths and B, Contained with the balance iron (Fe) and accompanying elements and impurities, mainly processed to be in the state of martensite phase with ultra-high strength and non-corrosive physical properties as a result of the choice and amount of ingredients and treatment protocol Including stainless steel alloys.

The subject invention alloy can achieve a maximum tensile strength (UTS) of about 300 ksi with a yield strength (YS) of about 230 ksi, and is about 6% and less than about 11% by weight, preferably about 10%. Also provides corrosion resistance over less than chromium. The alloys of the present invention provide a combination of the observed mechanical and corrosion properties of structural steels currently coated with cadmium and used in aerospace applications and without the special coating or plating. The highly efficient nanoscale carbide (M ₂ C) strengthening of the described alloys improves the corrosion resistance through the ability of nanoscale carbide to oxidize chromium and supply it as a passivating oxide, while reducing carbon and The alloy content provides ultra high strength. This combination of ultra high strength and corrosion resistance properties in a single material eliminates the need for a cadmium coating without the weight penalty associated with current structural steels. In addition, the alloys of the subject invention reduce wilderness failure driven by environmental embrittlement because they no longer rely on coatings that are unreliable for environmental protection.

  Accordingly, an object of the present invention is to provide a new class of ultra high strength corrosion resistant structural steel alloys.

  It is a further object of the present invention to provide an ultra high strength corrosion resistant structural steel alloy that does not require plating or coating to resist corrosion.

  Another object of the present invention is to provide an ultra high strength corrosion resistant structural steel alloy having cobalt, nickel and chromium alloy elements in combination with other elements, whereby the alloy is corrosion resistant.

  A further object of the invention is an ultra high strength having a maximum tensile strength (UTS) greater than about 240 ksi, preferably greater than about 280 ksi and yield strength (YS) greater than about 200 ksi, preferably greater than about 230 ksi. It is to provide a corrosion resistant structural steel alloy.

Another object of the present invention is characterized by a lath martensite microstructure and M ₂ C nanoscale size precipitation with a particle structure, and other M _x C precipitations where x> 2 are generally solubilized. It is to provide an ultra high strength corrosion resistant structural steel alloy.

  Yet another object of the present invention is to provide an ultra high strength corrosion resistant structural steel alloy that can be easily processed to form component parts and articles while maintaining ultra high strength and non-corrosive properties. is there.

  It is a further object of the present invention to provide a disclosed stainless steel alloy processing protocol that allows the creation of alloy microstructures with highly desirable strength and non-corrosion properties.

  These and other objects, advantages and features are described in the detailed description below.

(Detailed description of the invention)
In the following detailed description, reference is made to the drawings comprising the following figures: FIG.

  The steel alloy of the present invention exhibits various physical characteristics and throughput. These characteristics and capabilities were established as general standards, followed by the identification of suitable element combinations and processing steps to create such steel alloys and meet these standards. FIG. 1 is a system flow block diagram showing the processing / structure / property / performance relationship of an alloy of the present invention. Determine a set of alloy properties that require the desired performance for the application (eg, aerospace structure, landing gear, sports equipment, tools, etc.). The alloys of the present invention achieve the desired combination of properties and exhibit structural properties that can be evaluated through the continuous processing steps shown on the left of FIG. The following are criteria for alloy physical properties and throughput or characteristics. This is followed by a description of the analysis and experimental techniques associated with this invention, as well as examples of alloys that generally define the invention in terms of elemental regions and ranges, physical characteristics, and processing functions.

Physical Features The physical features or characteristics of the most preferred embodiments of the present invention are generally as follows:

Strength corresponding to or exceeding a corrosion-resistant 300M alloy corresponding to 15-5PH (H900 condition) as measured by linearly polarized light, ie

  Abrasion and fatigue resistance can be hardened to ≧ 67 Rockwell C (HRC).

  Microstructure features optimal for maximum fatigue / corrosion fatigue resistance.

Processability Features The main objective of the subject invention is to provide alloys with the physical properties listed above that have processability that makes them useful and practical. With some possible processing strategies related to the scale of manufacture for a given application and the resulting cleanliness and quality, compatibility of the subject invention alloys with a wide range of processes is desirable and is therefore the present invention. It is the feature.

  The main purpose and characteristics of this alloy are to perform melting, such as vacuum induction melting (VIM), vacuum arc remelting (VAR), and electroslag remelting (ESR) and other variations such as vacuum electroslag remelting. (VSR) compatibility. The alloys of the subject invention can also be made by other processes, such as air melting and powder metallurgy. What is important is the behavior of the alloy that exhibits limited solidification microsegregation under the solidification conditions of the above process. By selecting the appropriate elemental content in the alloys of the subject invention, compositional variations resulting from solidification across the secondary dendrites during processing can be minimized. Acceptable variation is a commercially viable temperature, usually a metal temperature above 1100 ° C. to the initial melting of the alloy, a reasonable processing time, typically less than 72 hours, preferably less than 36 hours. Yields an alloy that can be homogenized.

  The alloys of the subject invention also have reasonable hot ductility so that hot working after homogenization can be achieved within the temperature and reduction constraints typical of current industrial practice. Typical hot working for the alloys of the subject invention should allow a cross-section reduction ratio of greater than 3: 1 and preferably greater than 5: 1. In addition, initial hot working of the ingot should be possible below 1100 ° C. and final hot working to the desired product size should be possible at temperatures below 950 ° C.

For purposes of solution treatment, fine scale particle refining dispersions (ie MC) and generally ASTM particle size 5 by ASTM E112 [ASTM, ASTM El12-96, West Conshohocken, PA, 1996] incorporated herein. The goal is to completely dissolve all the main alloy carbides (ie, M _X C with X> 2) while maintaining a small particle size below. Thus, using the alloys of the present invention, during the solution treatment to the austenite phase region, the coarse scale alloy carbide formed during the pretreatment is dissolved and then the carbon in the resulting solution is tempered. It can be used for precipitation strengthening. However, during the same process, austenite particles can become coarse, thereby reducing strength, toughness and ductility. With the alloys of the present invention, such grain coarsening is slowed by MC precipitation that fixes grain boundaries and is necessary to avoid or reduce this grain coarsening as the solution treatment temperature increases. The amount of this particle refining dispersion increases. The alloys of the subject invention are all coarse scale carbides while maintaining an efficient particle refining dispersion at reasonable solution treatment temperatures in the range of 850 ° C. to 1100 ° C., preferably 950 ° C. to 1050 ° C. That is, M _x C where x> 2 is generally completely dissolved.

  After solution treatment, the components made from the subject invention alloy are typically rapidly cooled, i.e., quenched, below the temperature at which martensite is formed. The favorable result of this process consists essentially of all martensite and a microstructure that retains virtually no austenite, other transformation products such as bainite or ferrite or other carbide products remaining or formed during this process It is. The thickness of the component to be cooled and the cooling medium, such as oil, water, or air, determine the cooling rate for this type of process. As the cooling rate increases, the risk of other non-martensite product formation decreases, but the strain in the component potentially increases, thus reducing the cross-sectional thickness of the portion that can be processed. Alloys of the subject invention, when cooled to low temperature, or preferably to room temperature, are generally fully martensitic with a fragment size of less than 3 inches, preferably less than 6 inches, after cooling or quenching at a moderate rate. It is.

After cooling or quenching, the components produced using the subject invention alloys are cohesive to the carbon in the alloy while avoiding the formation of other carbide products, ie, M ₂ C where x> 2. It can be tempered at a temperature range and duration to form nanoscale M ₂ C carbide. During this aging or secondary curing process, the component is heated to the process temperature at a rate determined by the furnace power and the size of the component cross-section and held for a reasonable time before cooling or quenching to room temperature. Is done.

If pre-solution heat treatment is ineffective in avoiding austenite residues, this tempering process can be divided into multiple steps, where each tempering step is preferably cooled or quenched to room temperature and preferably Forms martensite through subsequent cooling to low temperatures. The temperature of this tempering process is typically 200 ° C. to 600 ° C., preferably 450 ° C. to 540 ° C., and the duration is less than 24 hours, preferably 2 to 10 hours. The desired process result is a nanoscale M ₂ C carbide dispersion that is free of transient cementite that forms early in the process and does not contain other alloy carbides that can precipitate if the processing time is excessively prolonged. It is a martensite matrix (generally austenite-free) reinforced by the object.

  An important feature of the alloys of the present invention relates to the high tempering temperatures used to achieve their secondary hardening response. A clear goal is to avoid cadmium plating for corrosion resistance, but many components made from the alloys of the present invention require an electroplating process, such as nickel or chromium, during manufacturing or overhaul It may be. The electroplating process introduces hydrogen into the microstructure, which can lead to embrittlement and must be baked by exposing the part to high temperatures after plating. The alloys of the present invention can be baked at temperatures as high as their original tempering temperatures without reducing the strength of the alloys. This bakeout process can be accomplished more quickly and reliably because the tempering temperatures are significantly higher in the alloys of the present invention compared to the commonly used 4340 and 300M alloys.

  Specific surface modification techniques for wear resistance, corrosion resistance, and decoration, such as physical vapor deposition (PVD), or surface hardening techniques, such as gas or plasma nitridation, at a temperature on the order of 500 ° C. This is done arbitrarily during the time order period. Another feature of the subject invention alloys is that the heat treatment process conforms to the temperature and schedule typical of these surface coating or curing processes.

  Components made of the subject invention alloys are typically manufactured or processed prior to solution treatment and aging. These manufacturing and processing operations require materials that are soft and exhibit the desired chip formation when the material is removed. Thus, the alloys of the subject invention are preferably annealed after the hot working process and before they are supplied to the manufacturer. The goal of this annealing process is to reduce the hardness of the subject invention alloy without promoting excessive austenite. Typically, annealing involves heating the alloy in the range of 600 ° C. to 850 ° C., preferably in the range of 700 ° C. to 750 ° C. for less than 24 hours, preferably 2 to 8 hours, and gradually cooling to room temperature. Achieved by: In some cases, a multi-step annealing process yields more optimal results. In such a process, the alloys of the present invention can be annealed at a series of temperatures for various times that may or may not be separated by an intermediate cooling step (s).

  After processing, solution treatment and aging, components made of the subject invention alloys may require a polishing step that maintains the desired final dimensions of the part. Surface polishing removes material from the part by the wear action on the high speed ceramic wheel. Damage to components due to heating of the surface of the part and damage to the grinding wheel due to material wear must be avoided. These troublesome problems can be avoided mainly by reducing the residual austenite content in the alloy. For this reason, and for other reasons discussed above, the subject invention alloys leave little austenite after solution treatment.

Many components made from the alloys of the subject invention may require joining by various welding processes, such as gas arc welding, underwater arc welding, friction stir welding, electron beam welding, and the like. These processes require a material that, after treatment, solidifies and becomes ductile in the melting or heating zone of the weld. Pre-heating and post-heating can be used to control the thermal history experienced by the alloy in the welding and heating zone and promote weld ductility. The main driving source for ductile welding is the lower carbon content in the material, but this also limits strength. The alloys of the subject invention achieve their strength using highly efficient nanoscale M ₂ C carbide and thus achieve a certain level of strength at a lower carbon content than steels such as 300M steel And consequently this promotes weldability.

Microstructural and compositional characteristics These alloy designs achieve the required corrosion resistance with minimal Cr content because the high Cr content limits other desirable properties in several ways. For example, one of the results of the high-Cr is reduced martensite M _s temperature, which in turn, other desirable alloying elements, for example, to limit the content and Ni. High Cr concentration also promotes excessive solidification microsegregation, which is difficult to eliminate by high temperature homogenization. High Cr also limits the high temperature solubilization of C required for carbide precipitation strengthening, which leads to the use of high solution treatment temperatures that make particle size control difficult. Thus, the characteristics of the alloys of the present invention are described as achieving corrosion resistance with structural strength of Cr in the range of greater than about 6% by weight and less than about 11% (preferably less than about 10%). In combination with other elements.

Another feature of the alloy is to achieve the required carbide strengthening with a minimum carbon content. Like the cr, C significantly reduced the M _s temperature, and raising the solution temperature. High C content also limits weldability and can cause corrosion problems associated with Cr carbide precipitation at grain boundaries. High C also limits the degree of softening that can be achieved by annealing to enhance machinability.

  Both of the main features just discussed are enhanced by the use of Co. The thermodynamic interaction of Co and Cr enhances the distribution of Cr to the oxide film formed during corrosion passivation and thus provides corrosion protection comparable to higher Cr steels. Co causes carbide precipitation during tempering by increasing the precipitation thermodynamic driving force and by delaying dislocation recovery to promote dissimilar nucleation of carbide by dislocation. Accordingly, C in the range of about 0.1% to 0.3% by weight combined with Co in the range of about 8% to 17% by weight, with Cr and other trace elements as described. Provides an alloy with corrosion resistance and ultra high strength.

The desired combination of corrosion resistance and ultra-high strength can also be improved by reducing the fineness of the carbide-reinforced dispersion to the nanostructure level, i.e., less than about 10 nanometers in diameter, preferably less than about 5 nanometers. . Compared to other strengthening precipitates, such as intermetallic phases used in maraging steel, the relatively high shear rate of M ₂ C alloy carbide reduces the optimal particle size for strengthening to a diameter of only about 3 nanometers Let Refinement to a carbide precipitation size to this level provides a very efficient reinforced dispersion. This is achieved by obtaining a sufficiently high thermodynamic driving force by alloying. This miniaturization brings the further advantage that the carbide has the same length scale as the passive oxide film, and the Cr in the carbide can participate in film formation. Therefore, the formation of carbide does not significantly reduce the corrosion resistance. A further advantage of nanoscale carbide dispersions is the hydrogen capture at the carbide interface that is effective in enhancing stress corrosion cracking resistance. Efficient nanoscale carbide strengthening is achieved by nitriding during tempering to make the system the same size scale M ₂ (C, N) carbonitriding for further efficient strengthening without significant loss of corrosion resistance Product should be produced and adequately suitable for surface hardening. Such nitriding can achieve a surface hardness as high as 1100 Vickers hardness (VHN) corresponding to 70 HRC.

Toughness is further enhanced by particle refining by optimal dispersion of particle refined MC carbide dispersion that maintains particle fixation during normalization and solution heat treatment and resists microvoid nucleation during ductile fracture . Melt deoxygenation practices are controlled to aid in the formation of Ti-rich MC dispersions for this purpose, as well as the number density of oxide and oxysulfide inclusion particles that form primary voids during fracture Is controlled to minimize. Under optimum conditions, the amount of MC determined by the mass balance from the available Ti content accounts for less than 10% of the alloy C content. Increasing Ni content within other demand constraints enhances resistance to brittle fracture. The refinement of M ₂ C particles by controlling the precipitation driving force is extremely high in order to completely dissolve Fe ₃ C cementite carbide that precipitates prior to M ₂ C and limits fracture toughness by forming microvoid nuclei. Allows strength to be maintained upon completion of M ₂ C precipitation. This cementite dissolution is described in Montgomery [Montgomery, J. et al. S. and G. B. Olson, M ₂ C Carbide Precipitation in AF1410, Gilbert R. Speich Symposium: Fundamentals of Aging and Tempering in Bainitic and Martensitic Steel Products, ISS-AIME, Warrendale, PA, 177-214, as evaluated by _{M 2} C Aie content measured using the techniques described by 1992, It is considered to be effectively completed when M ₂ C accounts for 85% of the alloy C content. Other phases that can limit toughness, such as other carbides (eg, M ₂₃ C ₆ , M ₆ C and M ₇ C ₃ ) and topological close packed (TCP) intermetallic phases (eg, σ and μ phases) ) Precipitation is avoided by constraining the thermodynamic driving force of their formation.

In addition to efficient hydrogen capture by nanoscale M ₂ C carbide that slows hydrogen transport, hydrogen stress is prevented by controlling the segregation of impurities and alloying elements to the pre-austenite grain boundary to prevent hydrogen assisted intergranular fracture. The resistance to corrosion is further enhanced. This controls the content of undesirable impurities, such as P and S, to low levels and getters their residual amount in the alloy to a stable compound, such as La ₂ O ₂ S or Ce ₂ O ₂ S To be promoted. Boundary aggregation is further enhanced by taking into account the segregation of aggregation strengthening elements such as B, Mo and W during heat treatment. These factors that promote stress corrosion cracking resistance also enhance resistance to corrosion fatigue.

All of these conditions are achieved while maintaining solution treatment temperatures that are not too high depending on the class of alloys discovered. Martensite M _s temperature measured by quenching expansion measurement and 1% transformation fraction also establish a lath martensite microstructure and minimize the content of residual austenite that can limit yield strength of the other way Maintained at a sufficient height.

Preferred processing techniques These alloys can be produced by various processes such as casting, powder metallurgy or ingot metallurgy. The alloy components can be melted by any conventional melting process such as air melting, but more preferably by vacuum induction melting (VIM). The alloy can then be homogenized and hot-processed, but to achieve improved fracture toughness and fatigue properties, secondary melting processes such as electroslag remelting (ESR) or vacuum arc remelting (VAR) Is preferred. Further remelting operations can be used prior to homogenization and hot working to achieve higher fracture toughness and fatigue properties. In any case, the alloy is first formed in the melting process by a combination of components.

  The alloy can then be homogenized prior to hot working or can be directly hot worked by heating. If homogenization is used, the alloy is soluble by heating to a metal temperature in the range of about 1100 ° C or 1110 ° C or 1120 ° C to 1330 ° C or 1340 ° C or 1350 ° C or possibly as high as 1400 ° C for a period of at least 4 hours. This can be done by dissolving the elements and carbides and making the structure uniform. One of the alloy design criteria is low microsegregation, and thus the time required to homogenize the alloy is typically shorter than other stainless steel alloys. A suitable time is 6 hours or more in the homogenized metal temperature range. In general, the soaking time at the homogenization time should not exceed 72 hours. It has been found that 12 to 18 hours is very suitable in the homogenization temperature range. A typical homogenizing metal temperature is about 1240 ° C.

  After homogenization, the alloy is typically hot worked. The alloy can be hot worked by, but not limited to, hot rolling, hot forging or hot extrusion or any combination thereof. In order to utilize the heat already present in the alloy, it is common to start hot working immediately after the homogenization process. It is important that the final hot-working metal temperature is substantially below the starting hot-working metal temperature in order to ensure grain refinement of the structure due to MC carbide precipitation. After the first hot working step, the alloy is typically reheated to continue hot working to the final desired size and shape. The reheat metal temperature range is about 950 ° C. or 960 ° C. or 970 ° C. to 1230 ° C. or 1240 ° C. or 1250 ° C. or possibly as high as 1300 ° C., with a preferred range of about 1000 ° C. or 1010 ° C. to 1150 ° C. Or it is 1160 degreeC. The reheat metal temperature is near or approximately the MC carbide solvus temperature, the purpose of which is to dissolve or partially dissolve soluble components that may remain from the casting or precipitate during prior hot working. To dissolve. This reheating process minimizes or avoids primary and secondary phase particles and improves fatigue crack growth resistance and fracture toughness.

  Because the alloy is continuously hot worked and reheated, the cross-sectional size is reduced and, as a result, the metal cools more rapidly. Eventually, high reheat temperatures can no longer be used and lower reheat temperatures must be used. For smaller cross sections, the reheat metal temperature range is about 840 ° C or 850 ° C or 860 ° C to 1080 ° C or 1090 ° C or 1100 ° C or possibly as high as 1200 ° C, with a preferred range of about 950 ° C. Or it is 960 degreeC to 1000 degreeC or 1010 degreeC. The lower reheat metal temperature for smaller cross-sections is below the solvus temperature of other (non-MC) carbides, the purpose of which is to quickly dissolve during the next normalization or solution treatment As is, to minimize or prevent their coarsening during reheating.

  Final mill product forms, such as bar stock and forged steel stock, are typically tempered and / or tempered before shipping to the consumer. During normalization, the alloy is heated to a metal temperature above the solvus temperature of all carbides except MC carbide, the purpose of which is to dissolve the soluble components that can precipitate during the hot working prior to it and to reduce the particle size. The normalizing metal temperature range to be leveled is about 880 ° C or 890 ° C or 900 ° C to 1080 ° C or 1090 ° C or 1100 ° C, with a preferred range being about 1020 ° C to 1030 ° C or 1040 ° C. A suitable time is 1 hour or more and typically the soaking time at the normalizing temperature should not exceed 3 hours. Thereafter, the alloy is cooled to room temperature.

  After normalization, the alloy is typically annealed to a hardness or strength level suitable for subsequent consumer processing, eg, machining. During annealing, the alloy is heated to a metal temperature range of about 600 ° C. or 610 ° C. to 840 ° C. or 850 ° C., preferably 700 ° C. to 750 ° C. for a period of at least 1 hour to coarsen all the carbides except MC carbide To do. A suitable time is 2 hours or more, and typically the soaking time at the annealing temperature should not extend beyond 24 hours.

  Typically, after the alloy is delivered to the consumer and processed to or near its final form, it is preferably about 850 ° C. or 860 ° C. to 1090 ° C. or 1100 ° C., more preferably about 950 ° C. To solution treatment for a period of less than 3 hours at a metal temperature range of 1040 ° C or 1050 ° C. A typical time for solution treatment is one hour. This solution treated metal temperature is above the solvus temperature of all carbides except MC carbide, the purpose of which is to dissolve soluble components that may precipitate during the preceding process. This prevents grain growth while enhancing strength, fracture toughness and fatigue resistance.

  After solution treatment, it is important to cool the alloy to room temperature quickly enough to transform the microstructure primarily into a lath martensite structure and prevent or minimize the boundary precipitation of the main carbide. is there. A suitable cooling rate can be achieved by using water, oil, or various quenching gases, depending on the cross-sectional thickness.

  After quenching to room temperature, the alloy can be subjected to a low temperature treatment or directly heated to the tempering temperature. Low temperature treatment promotes a more complete transformation to the microstructure lath martensite structure. When low temperature processing is used, it is preferably performed at less than about -70 ° C. A more preferred low temperature treatment is less than about -195 ° C. A typical low temperature treatment is a metal temperature range of about -60 ° C or -70 ° C to -85 ° C or -95 ° C. Another typical low temperature treatment is a metal temperature range of about -180 ° C or -190 ° C to 220 ° C or -230 ° C. Normally, the soaking time at low temperatures should not exceed 10 hours. A typical time for low temperature treatment is 1 hour.

The alloy is tempered at the intermediate metal temperature after the low temperature treatment or immediately after quenching if the low temperature treatment is omitted. This tempering treatment is preferably in the metal temperature range of about 200 ° C or 210 ° C or 220 ° C to 580 ° C or 590 ° C or 600 ° C, more preferably about 450 ° C to 530 ° C or 540 ° C. Usually, the soaking time at the tempering temperature should not exceed 24 hours. It has been found that 2 to 10 hours are very suitable in this tempering temperature range. During this tempering process, the precipitation of nanoscale M ₂ C strengthened particles increases the thermal stability of the alloy and various combinations of strength and fracture toughness can be achieved by using different combinations of temperature and time. .

  For alloys of the present invention with lower MS temperatures, it is possible to further strengthen strength and fracture toughness by minimizing retained austenite with a multi-step heat treatment. Multi-step processing consists of an additional cycle of low-temperature processing and subsequent heat treatment outlined in the text above. While one additional cycle may be beneficial, typically multiple cycles are more beneficial.

  An example of the relationship between processing strategy and phase stability in a particular alloy of the present invention is shown in FIGS. 2A and 2B.

  FIG. 2A shows the equilibrium phase of Alloy 2C of the present invention having a carbon content of 0.23% by weight as shown in Table 1.

  FIG. 2B discloses the processing sequence used for the described alloy 2C. After forming the melt by the melt processing step, the alloy is homogenized at a metal temperature above the single phase (fcc) equilibrium temperature of about 1220 ° C. All carbides are solubilized at this temperature. Forging defining the desired billet, rod or other shape results in cooling to the extent that various composite carbides can be formed. This forging process can be repeated at least by reheating to a metal temperature range (980 ° C. to 1220 ° C.) where only MC carbide is in equilibrium.

Subsequent cooling (air cooling) generally retains mainly the formation of MC carbide, other main alloy carbides such as M ₇ C ₃ and M ₂₃ C ₆ and generally the martensite matrix. Normalization and subsequent cooling in the same metal temperature range dissolves the M ₇ C ₃ and M ₂₃ C ₆ primary carbides while preserving the MC carbides. Annealing and cooling in the metal temperature range of 600 ° C. or 610 ° C. to 840 ° C. or 850 ° C. reduces the hardness level to a reasonable value for machining. The annealing process softens martensite by precipitating carbon in alloy carbide that is too large to significantly strengthen the alloy but is small enough to be easily dissolved during a subsequent solution heat treatment. This process is followed by final manufacture of the component parts and shipment of the alloy product to the consumer for appropriate heat treatment and finishing.

  Typically, the consumer forms the alloy into the desired shape. This is followed by a solution treatment in the MC carbide temperature range, followed by quenching to maintain or form the desired martensite structure. The tempering and cooling described above can then be used to obtain the desired strength and fracture toughness.

Experimental results and examples A series of prototype alloys were prepared. The melting operation for the miniaturization process was chosen to be double vacuum melting with La and Ce impurity gettering addition. Substitute particle boundary aggregation enhancers such as W and Re were not considered in making the first protocol, but the addition of 20 ppm B was included for this purpose. Because of the deoxygenation protocol, Ti was added as a deoxygenating agent prior to tempering to promote TiC particles to fix the particle boundaries and reduce particle growth during the solution heat treatment.

The main alloying elements in the first prototype are C, Mo, and V (M ₂ C carbide formers), Cr (M ₂ C carbide formers and oxide passive film formers), and Co and Ni (various requirements) For matrix properties). The actual alloy composition and material processing parameters are incorporated into the overall design integration considering chaining and a set of computer models described elsewhere [Olson, G. et al. B, “Computational Design of Hierarchically Structured Materials.”, Science 277, 1237-1242, 1997]. The following is a summary of the initial prototype procedure. The selected parameter is a star in Figure 3-6.

Indicated by.

The amount of Cr is the corrosion resistance requirement and a passive thermodynamic model developed by Cambell, incorporated herein [Campbell, C, Systems Design of High Performance Stainless Steels, 19 Materials Eight, 19th, Environs Science and Eng. ]. The amount of C was determined by an M2C precipitation / strengthening model with strength requirements and the correlation shown in FIG. Based on the goal of achieving 53 HRC hardness, a C content of 0.14% by weight was selected. The tempering temperature and the amounts of M2C carbide formers Mo and V were determined to meet strength requirements with appropriate M2C precipitation kinetics, maintain a solution heat treatment temperature of 100 ° C., and avoid microsegregation. Figures 4 and 5 show how the final V and Mo contents were determined. A final content of 1.5 wt% Mo and 0.5 wt% V was selected. The level of solidification microsegregation is assessed by solidification simulation for the expected ingot solidification cooling rate and associated dendrite arm spacing. The amount of Co and Ni is maintained at a martensite onset temperature of at least 200 ° C. using a model calibrated to (1) quench expansion measurements and 1% transformation fraction, and therefore lath after quenching. A martensitic matrix structure can be achieved, (2) high M2C carbide initial driving force for efficient strengthening is maintained, (3) bcc cleavage resistance is improved by maximizing Ni content, and (4) Decided to maintain a Co content of about 8% by weight to achieve dislocation recovery resistance sufficient to enhance M ₂ C nucleation, and to increase Cr distribution to the oxide film due to increased matrix Cr activity To do. FIG. 6 shows that optimization of the above four factors, along with other alloy element amounts and tempering temperatures set at their final levels, resulted in Co and Ni amounts of about 13% and 4.8% by weight, respectively. Indicates that a selection will occur. These material composition and tempering temperature refers to past studies of other precipitation hardened Ni-Co steels were fine-tuned by inspecting the driving force ratios between the M 2 _C and other carbides and intermetallic phases .

  The composition of the first design prototype, designated 1, is shown in Table 1 along with subsequent design iterations. The initial design included the following processing parameters:

-Double vacuum melting with impurity gettering and Ti deoxygenation,
• Minimum solution heat treatment temperature of 1005 ° C., limited according to thermodynamic equilibrium by vanadium carbide (VC) formation, and • Tempering at 482 ° C. with an estimated tempering time of 3 hours to achieve optimum strength and toughness temperature.

Evaluation of the first prototype (entry 1 in Table 1) gave promising results for all properties evaluated. The most serious disadvantages were 25-50 ° C. below the desired _Ms temperature and an intensity level 15% below the target. Next, a second series of designs designated 2A, 2B and 2C in Table 1 was evaluated. All three second iteration prototypes provided satisfactory transformation temperatures, and the best mechanical properties of the second iteration were demonstrated by Alloy 2C. Based on the latter basic composition, the third iterative series of alloys, referred to in Table 1 as 3A, 3B and 3C, compared the TiC, (Ti, V) C, and NbC, with small variations in grain refined MC carbide Explored. The main parameters were the MC phase fraction at solution temperature and the resistance to coarsening, affected by the constraint of complete MC solubility at the homogenization temperature. With (Ti, V) C selected as the optimal grain refinement approach, the fourth iterative design series, referred to as 4A to 4G in Table 1, is (a) a refinement of martensitic transformation kinetics that minimizes retained austenite content. (B) Promote complete dissolution of cementite during M ₂ C precipitation strengthening to enhance fracture toughness, increase competitive M ₂ C carbide stability, and (c) completely avoid cementite precipitation And low temperature iron (Fe) based M ₂ C precipitation strengthening used to increase cleavage resistance. Changes in carbide thermodynamics and kinetics in the latter two series included the addition of W and Si.

  The fifth series of alloys, referred to as 5B to 5F in Table 1, tested the limit of Ni that can be added to the alloy to improve fracture toughness by reducing the transition temperature from ductility to brittleness. When the Ni content reaches about 10 percent by weight, tempering of the alloy in multiple steps with low temperature cooling between each step keeps most of the retained austenite martensite while keeping the alloys Ms of these compositions below room temperature. It was found that it can be converted to a site. This makes it possible to control ductile fracture behavior by achieving good strength properties in combination with high Ni content even in alloys that are fully austenitic after quenching. Although multiple tempers are commonly used to minimize residual austenite in steel, it is unexpected that the technique can be used effectively in alloys with such high Ni and high austenite contents. It was.

  The sixth series of alloys, referred to as 6A to 6M in Table 1, were determined to incorporate the features represented in the first five series and are considered preferred aspects of the present invention. Thus, proper processing of the described alloy provides an essentially martensitic phase.

The following is a summary of the experiments and alloys described.

(Example 1)
Alloy 1 in Table 1 was vacuum induction melted (VIM) on a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) on an 8 inch diameter ingot. The material was homogenized at 1200 ° C. for 72 hours, forged and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 175 ° C. by quenching expansion measurement and 1% transformation fraction.

The test sample was machined, solution treated at 1025 ° C. for 1 hour, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 482 ° C. for 8 hours. The measured properties are listed in Table 2 below.

(Example 2)
Alloy 2A in Table 1 was vacuum induction melted (VIM) on a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) on an 8 inch diameter ingot. The ingot was homogenized for 12 hours at 1190 ° C. for 12 hours, forged and rolled into 1.500 inch square bar steel and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 265 ° C. by quenching expansion measurement and 1% transformation fraction.

Test samples are machined from square bar steel, solution treated at 1050 ° C. for 1 hour, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 500 ° C. for 5 hours, air cooled, liquid nitrogen For 1 hour, warmed to room temperature, and tempered at 500 ° C. for 5.5 hours. The measured properties are listed in Table 3 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.26 mil (mpy) annually.

Tensile samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and 1 hour of liquid nitrogen (LN ₂₎ during the tempering process. ) And multi-step tempering at 496 ° C. for 4 or 6 hours. The measured tensile properties are listed in Table 4 below.

(Example 3)
Alloy 2B in Table 1 was vacuum dielectric melted (VIM) into a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) into an 8 inch diameter ingot. The ingot was homogenized for 12 hours at 1190 ° C. for 12 hours, forged and rolled into a 1.000 inch round steel bar and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 225 ° C. by quenching expansion measurement and 1% transformation fraction.

Test samples were machined from round bar steel, solution treated at 1100 ° C. for 70 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 482 ° C. for 24 hours. The measured properties are listed in Table 5 below.

(Example 4)
Alloy 2C in Table 1 was vacuum dielectric melted (VIM) into a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) into an 8 inch diameter ingot. The ingot was homogenized at 1190 ° C. for 12 hours, forged into a 2.250 inch square bar steel, starting at 1120 ° C. and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 253 ° C. by quenching expansion measurement and 1% transformation fraction.

The test sample was machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 498 ° C. for 8 hours. The measured properties are listed in Table 6 below.

Test samples were machined from square bar steel, solution treated at 1025 ° C for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 498 ° C for 12 hours. The measured properties are listed in Table 7 below.

Corrosion test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 498 ° C. for 8 hours, air cooled, 498 Tempering at 4 ° C. for 4 hours. The measured properties are listed in Table 8 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.26 mil (mpy) annually.

Tensile samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and 1 hour of liquid nitrogen (LN ₂₎ during the tempering process. ) And multi-step tempering at 496 ° C. for 4 or 6 hours. The measured tensile properties are listed in Table 9 below.

(Example 5)
Alloy 3A in Table 1 was vacuum dielectric melted (VIM) into a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) into an 8 inch diameter ingot. The ingot was homogenized at 1260 ° C. for 12 hours, forged into 2.250 inch square bar steel, starting at 1090 ° C. and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 250 ° C. by quenching expansion measurement and 1% transformation fraction.

Test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 5 hours. The measured properties are listed in Table 10 below.

Test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 510 ° C. for 4 hours and multi-step tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 510 ° C. for an additional 4 hours. The measured properties are listed in Table 11 below.

The corrosion test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 200 ° C. for 1 hour. The measured properties are listed in Table 12 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.20 mil (mpy) annually.

The corrosion test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 8 hours. The measured properties are listed in Table 13 below.

(Example 6)
Alloy 3B in Table 1 was vacuum dielectric melted (VIM) into a 6 inch diameter electrode and then it was vacuum arc remelted (VAR) into an 8 inch diameter ingot. The ingot was homogenized at 1260 ° C. for 12 hours, forged into 2.250 inch square bar steel, starting at 1090 ° C. and annealed according to the preferred processing technique described above and shown in FIGS. 2A and 2B. Dilatometer samples were machined, M _s temperature was measured as 240 ° C. by quenching expansion measurement and 1% transformation fraction.

Test samples were machined from square bar steel, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and finally tempered at 510 ° C. for 5 hours. The measured properties are listed in Table 14 below.

The test sample was machined from a square bar, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 510 ° C. for 4 hours and then tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 510 ° C. for an additional 4 hours. The measured properties are listed in Table 15 below.

(Example 7)
Alloy 4A in Table 1 was vacuum dielectric melted (VIM) on a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) on a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 275 ° C. by quenching expansion measurement and 1% transformation fraction.

Corrosion test samples were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 12 hours. The measured properties are listed in Table 16 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.20 mil (mpy) annually.

Corrosion test samples were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 24 hours. The measured properties are listed in Table 17 below.

(Example 8)
Alloy 4B in Table 1 was vacuum dielectric melted (VIM) into a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) into a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 285 ° C. by quenching expansion measurement and 1% transformation fraction.

The corrosion test sample was machined from a rectangular bar, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 12 hours. The measured properties are listed in Table 18 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.20 mil (mpy) annually.

Example 9
Alloy 4C in Table 1 was vacuum dielectric melted (VIM) on 4 inch diameter electrodes, then it was vacuum arc remelted (VAR) on 5 inch diameter ingots. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 310 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 200 ° C. for 2 hours and then tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 200 ° C. for another 2 hours. The measured properties are listed in Table 19 below.

(Example 10)
Alloy 4D in Table 1 was vacuum dielectric melted (VIM) into a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) into a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 300 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 200 ° C. for 2 hours and tempered for 2 hours. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 200 ° C. for another 2 hours. The measured properties are listed in Table 20 below.

Corrosion test samples were machined from rectangular steel bars, solution treated at 1000 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 12 hours. The measured properties are listed in Table 21 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.20 mil (mpy) annually.

(Example 11)
Alloy 4E in Table 1 was vacuum dielectric melted (VIM) on a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) on a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 300 ° C. by quenching expansion measurement and 1% transformation fraction.

(Example 12)
Alloy 4F in Table 1 was vacuum dielectric melted (VIM) on 4 inch diameter electrodes, then it was vacuum arc remelted (VAR) on 5 inch diameter ingots. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 300 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 200 ° C. for 2 hours and tempered for 2 hours. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 200 ° C. for another 2 hours. The measured properties are listed in Table 22 below.

Corrosion test samples were machined from rectangular steel bars, solution treated at 1000 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and tempered at 510 ° C. for 12 hours. The measured properties are listed in Table 23 below. Reference to the corrosion rate of 15-5PH (H900 condition) was made using samples tested under the same conditions. The average corrosion rate of 15-5PH (H900 condition) for this test was 0.20 mil (mpy) annually.

(Example 13)
Alloy 4G in Table 1 was vacuum dielectric melted (VIM) on 4 inch diameter electrodes, then it was vacuum arc remelted (VAR) on 5 inch diameter ingots. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 320 ° C. by quenching expansion measurement and 1% transformation fraction.

(Example 14)
Alloy 5B in Table 1 was vacuum dielectric melted (VIM) on a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) on a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 200 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 24 hours and then tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 24 hours. The measured properties are listed in Table 24 below.

The test sample was machined from a rectangular bar, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 36 hours, and tempered for 1 hour. Liquid nitrogen (LN ₂ ) treatment and finally tempering at 468 ° C. for a further 36 hours. The measured properties are listed in Table 25 below.

(Example 15)
Alloy 5C in Table 1 was vacuum dielectric melted (VIM) into a 4 inch diameter electrode and then it was vacuum arc remelted (VAR) into a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 180 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 16 hours and tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 16 hours. The measured properties are listed in Table 26 below.

(Example 16)
Alloy 5D in Table 1 was vacuum dielectric melted (VIM) on a 4 inch diameter electrode, then it was vacuum arc remelted (VAR) on a 5 inch diameter ingot. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 240 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 24 hours and then tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 24 hours. The measured properties are listed in Table 27 below.

Test samples were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and finally tempered at 468 ° C. for 28 hours. The measured properties are listed in Table 28 below.

Test samples were machined from rectangular steel bars, solution treated at 1025 ° C for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and finally tempered at 468 ° C for 72 hours. The measured properties are listed in Table 29 below.

(Example 17)
Alloy 5E in Table 1 was vacuum dielectric melted (VIM) on 4 inch diameter electrodes, then it was vacuum arc remelted (VAR) on 5 inch diameter ingots. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured as 165 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 16 hours and tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 16 hours. The measured properties are listed in Table 30 below.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 24 hours and then tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 24 hours. The measured properties are listed in Table 31 below.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 14 hours and multi-step tempered for 1 hour. Liquid nitrogen (LN ₂ ) treatment and finally tempering at 468 ° C. for 14 hours. The measured properties are listed in Table 32 below.

(Example 18)
Alloy 5F in Table 1 was vacuum dielectric melted (VIM) on 4 inch diameter electrodes, then it was vacuum arc remelted (VAR) on 5 inch diameter ingots. The ingot was homogenized for 12 hours at 1250 ° C. for 12 hours according to the preferred processing technique described above and shown in FIGS. 2A and 2B, and hot rolled into a 2 inch square round using frequent reheating at 1015 ° C. Hot rolled into a 750 inch thick by 2.250 inch wide rectangular steel bar, normalized and annealed. Dilatometer samples were machined, M _s temperature was measured less than 25 ° C. by quenching expansion measurement and 1% transformation fraction.

Test specimens were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, tempered at 468 ° C. for 16 hours and tempered for 1 hour. It was treated with liquid nitrogen (LN ₂ ) and finally tempered at 468 ° C. for an additional 16 hours. The measured properties are listed in Table 33 below.

Test samples were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and finally tempered at 468 ° C. for 28 hours. The measured properties are listed in Table 34 below.

Test samples were machined from rectangular steel bars, solution treated at 1025 ° C. for 75 minutes, oil cooled, immersed in liquid nitrogen for 1 hour, warmed to room temperature, and finally tempered at 468 ° C. for 48 hours. The measured properties are listed in Table 35 below.

  The examples in Tables 34 and 35 show the advantages of multi-step tempering of alloys that provide higher strength.

The key to alloy design is to achieve efficient strengthening while maintaining effective hydrogen capture for corrosion resistance and stress corrosion resistance. All of these properties are facilitated by refining the toughening M ₂ C carbide particle size to an optimal size of about 3 nanometers upon completion of deposition. FIG. 7 shows atomic scale imaging of 3 nanometer M ₂ C carbide in optimally heat-treated alloy 2C using 3D atom-probe microanalysis and is incorporated herein [M. K. Miller, Atom Probe Tomography, Kluwer Academic / Plenum Publishers, New York, NY, 2000] demonstrates that the designed size and particle composition have actually been achieved. This image is an atomic reconstruction of the alloy slab, where each atom is represented as a point on the diagram having a color and size corresponding to that element. The circle depicted in FIG. 7 represents a collection of alloy carbide formers and carbon defining M ₂ C nanoscale carbide in the image.

As a result, the discovered alloy has a series of combinations of the elements shown in Table 36.

Preferably, impurities are avoided, however, some impurities and accidental elements are tolerated and within the scope of the present invention. Thus, by weight, most preferably, S is less than 0.02%, P is less than 0.012%, O is less than 0.015% and N is less than 0.015%. The microstructure is primarily martensitic when processed as described, desirably maintained as lath martensitic with less than 2.5% by volume, preferably less than 1% residual or precipitated austenite. Is done. The microstructure mainly comprises M ₂ C nanoscale carbide, where M is one or more elements selected from the group comprising Mo, Nb, V, Ta, W and Cr. Carbide formula, size and presence are important. Preferably, the carbide is in the form of M ₂ C only, and to some extent MC carbide, present in the absence of other carbides, and its size (average diameter) is less than about 10 nanometers, preferably about 3 nanometers. It is in the range of 5 to 5 nanometers. Particularly avoided are other larger incoherent carbides such as cementite, M ₂₃ C ₆ , M ₆ C and M ₇ C ₃ . Other embrittlement phases are also avoided, for example, topological close packed (TCP) intermetallic phases.

The martensite matrix in which the reinforced nanocarbide is embedded is high enough to enhance Cr distribution to the passivating oxide, enhance M ₂ C driving force, and maintain dislocation nucleation of the nanocarbide. In order to maintain the M _s temperature with sufficient Co, an optimal balance of Co and Ni is included. Cleavage resistance is enhanced by promoting sufficient grain refinement with a stable MC carbide dispersion that maintains sufficient Ni and resists coarsening at normalizing or solution heat treatment temperatures. The alloy composition and heat treatment are optimized to minimize or eliminate all other dispersed particles that limit toughness and fatigue resistance. Resistance to hydrogen stress corrosion is enhanced by grain boundary segregation of cohesive strengthening elements such as B, Mo and W, and by the hydrogen trapping effect of nanoscale M ₂ C carbide dispersions. The alloy composition is constrained to limit fine segregation under product-scale ingot solidification conditions.

The specific alloy composition in Table 1 represents the currently known preferred and optimal formulations for this class of alloys, with formulation variations consistent with the physical properties described, process variations and variations within the disclosed scope. Other equivalents are within the scope of the invention. These preferred embodiments can be summarized as seven subclasses of alloy compositions shown in Table 37. Subclass 1 is similar to alloys 2C, 3A and 3B in Table 1 and has a secondary hardening of about 400 ° C. to 600 ° C. which precipitates Cr—Mo based M ₂ C carbide resulting in UTS in the range of about 270 ksi to 300 ksi. Ideal for temperature. Subclass 2 is similar in composition to Alloys 4D and 4E in Table 1, with a secondary hardening temperature of about 400 ° C. to 600 ° C. that destabilizes cementite and precipitates Cr—Mo—W based M ₂ C carbide. Includes the addition of W and / or Si that provides high thermal stability. For applications that require higher fracture toughness, subclass 3 is similar in composition to alloys 1, 2A and 2B in Table 1, resulting in an intermediate UTS range of about 240 ksi to 270 ksi. Subclass 4 is similar in composition to Alloys 4F and 4G in Table 1 and is optimal for low temperature tempering at about 200 ° C. to 300 ° C. to precipitate Fe-based M ₂ C carbide without precipitation of cementite. Alloy subclass 5 is the most preferred embodiment of subclass 1. Subclass 6 is similar in composition to alloys 5B to 5F and 6A to 6K. Subclass 6 provides optimal toughness due to the high Ni content, but may require multiple tempering treatments with low temperature treatments between steps to avoid large amounts of retained austenite in the final microstructure. Subclass 7 is a further optimization of fracture toughness, similar to alloys 6L and 6M, where low Co content reduces the ductile-brittle transition temperature of the alloy.

  Accordingly, the present invention, including the class of ultra high strength corrosion resistant structural steel alloys and methods of making and using such alloys, is limited only by the following claims and their equivalents.

FIG. 3 is a flow block logic diagram characterizing the design concept of the alloy of the present invention. FIG. 2A is an equilibrium phase diagram illustrating the carbide phase and composition at various temperatures in an example of an alloy of the present invention. FIG. 2B is a diagram of an exemplary processing strategy for an alloy of the present invention with respect to the proposed equilibrium phase. FIG. 4 is a graph correlating hardness peaks and M ₂ C driving force at weight percent values for various carbon (C) contents. FIG. 5 is a graph showing the gradient of M ₂ C driving force (ΔG) and the scale rate constants for various molybdenum (Mo) and vanadium (V) contents, where the temperature is set to 482 ° C. and other alloy elements Is set to 0.14 wt% carbon (C), 9 wt% chromium (Cr), 13 wt% cobalt (Co), and 4.8 wt% nickel (Ni). FIG. 2 is a phase diagram at 1000 ° C. used to determine the final vanadium (V) content for a carbon (C) content of 0.14 wt%, where the amount of other alloying elements is 9 wt% chromium (Cr). 1.5 wt% molybdenum (Mo), 13 wt% cobalt (Co), and 4.8 wt% nickel (Ni). FIG. 6 is a graph showing M _s temperature gradient and M ₂ C driving force (ΔG) for various cobalt (Co) and nickel (Ni) contents in one embodiment of the present invention, where the temperature is set to 482 ° C. And other alloy element amounts are set to 0.14 wt% carbon (C), 9 wt% chromium (Cr), 1.5 wt% molybdenum (Mo), and 0.5 wt% vanadium (V). Has been. 3 is a three-dimensional atom-probe image of M ₂ C carbide in an optimally heat-treated preferred embodiment and example of the present invention.

Claims (173)

  About 0.1 to 0.3% carbon (C), 8 to 17% cobalt (Co), less than 5% nickel (Ni), greater than 6 and less than 11% chromium (Cr) by weight And an alloy composition comprising less than 3% molybdenum (Mo), essentially the remainder of iron (Fe), and a combination of incidental elements and impurities.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 240 ksi.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 260 ksi.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 280 ksi.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 240 ksi and a yield strength (YS) greater than about 200 ksi.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 260 ksi and a yield strength (YS) greater than about 215 ksi.   128. The alloy of claim 1 or claim 127, having a maximum tensile strength (UTS) greater than about 280 ksi and a yield strength (YS) greater than about 230 ksi. 128. The alloy of claim 1 or claim 127 having a martensite onset (M _s ) temperature measured by a rapid expansion measurement greater than about 150 ° C. and a 1% transformation fraction. 128. The alloy of claim 1 or claim 127 having a martensite onset ( _Ms ) temperature measured by a rapid expansion measurement greater than about 200 <0> C and a 1% transformation fraction. 128. The alloy of claim 1 or claim 127 having a martensite onset (M _s ) temperature measured by a rapid expansion measurement greater than about 250 ° C. and a 1% transformation fraction. Alloys containing less than about 10 nanometers of M ₂ C carbide have a carbon (C) content greater than about 85% by weight, where M is Cr, Mo, V, W, Nb, Ta and their 128. The alloy of claim 1 or claim 127, selected from the group comprising combinations. Alloys containing less than about 5 nanometers of M ₂ C carbide have a carbon (C) content of greater than about 85% by weight, where M is Cr, Mo, V, W, Nb, Ta and their 128. The alloy of claim 1 or claim 127, selected from the group comprising combinations.   Equal to the rate determined for 15-5PH (H900 condition) stainless steel as measured by linear polarimetry in a 3.5% by weight aqueous sodium chloride solution formed with a Cr passivated surface layer 128. The alloy of claim 1 or claim 127 having an annual corrosion rate of or less.   About a rate determined for 15-5PH (H900 condition) stainless steel, as measured by linear polarimetry in a 3.5% by weight aqueous sodium chloride solution formed with a Cr passivated surface layer. 128. The alloy of claim 1 or claim 127 having an annual corrosion rate of less than 250%.   128. The alloy of claim 1 or claim 127 having a maximum tensile strength (UTS) greater than about 240 ksi and a martensite onset (Ms) temperature measured by a rapid expansion measurement and a 1% transformation fraction greater than about 200 ° C. .   128. The alloy of claim 1 or claim 127 having a maximum tensile strength (UTS) greater than about 260 ksi and a martensite onset (Ms) temperature measured by a rapid expansion measurement and a 1% transformation fraction greater than about 200 ° C. .   128. The alloy of claim 1 or claim 127 having a maximum tensile strength (UTS) greater than about 280 ksi and a martensite onset (Ms) temperature measured by a rapid expansion measurement and a 1% transformation fraction greater than about 200 ° C. .   A maximum tensile strength (UTS) greater than about 240 ksi and a rate of about 250 determined for 15-5 PH (H900 condition) stainless steel as measured by linear polarimetry in 3.5% by weight aqueous sodium chloride solution. 128. The alloy of claim 1 or claim 127 having an annual corrosion rate of less than%.   A maximum tensile strength (UTS) greater than about 260 ksi and a rate of about 250 determined for 15-5PH (H900 condition) stainless steel as measured by linear polarimetry in 3.5% by weight aqueous sodium chloride solution. 128. The alloy of claim 1 or claim 127 having an annual corrosion rate of less than%.   A maximum tensile strength (UTS) greater than about 280 ksi and a rate of about 250 determined for 15-5PH (H900 condition) stainless steel as measured by linear polarimetry in 3.5% by weight aqueous sodium chloride solution. 128. The alloy of claim 1 or claim 127 having an annual corrosion rate of less than%.   Is equal to the rate determined for 15-5PH (H900 condition) stainless steel as measured by linear polarimetry in a maximum tensile strength (UTS) greater than about 240 ksi and 3.5% by weight aqueous sodium chloride solution? 128. The alloy of claim 1 or claim 127, having an annual corrosion rate of, or less.   Is the maximum tensile strength (UTS) greater than about 260 ksi and the rate determined for 15-5 PH (H900 condition) stainless steel as measured by linear polarimetry in a 3.5% by weight aqueous sodium chloride solution? 128. The alloy of claim 1 or claim 127, having an annual corrosion rate of, or less.   Is the maximum tensile strength (UTS) greater than about 280 ksi and equal to the rate determined for 15-5 PH (H900 condition) stainless steel as measured by linear polarimetry in 3.5% by weight aqueous sodium chloride solution? 128. The alloy of claim 1 or claim 127, having an annual corrosion rate of, or less. Alloys having a maximum tensile strength (UTS) greater than about 240 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 10 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 240 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 5 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 260 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 10 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 260 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 5 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 280 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 10 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 280 ksi and greater than about 85% by weight of the carbon content are found in M ₂ C carbides less than about 5 nanometers, where M is Cr 128. The alloy of claim 1 or claim 127, selected from the group consisting of: Mo, V, W, Nb, Ta, and combinations thereof. Alloys having a maximum tensile strength (UTS) greater than about 240 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 10 nanometers, where M is Cr , Mo, V, W, Nb, Ta, and combinations thereof, the martensite onset (M _s ) temperature of the alloy measured by quench expansion measurement and 1% transformation fraction is greater than about 150 ° C. And the annual corrosion rate measured by linear polarimetry in a 3.5 wt% aqueous sodium chloride solution is less than about 250% of the rate determined for 15-5PH (H900 condition). Alloy. Alloys having a maximum tensile strength (UTS) greater than about 240 ksi and greater than about 85% by weight are found in M ₂ C carbides where the carbon content is less than about 5 nanometers, where M is Cr , Mo, V, W, Nb, Ta and combinations thereof, the martensite onset (M _s ) temperature of the alloy measured by quench expansion measurement and 1% transformation fraction is greater than about 150 ° C. And the annual corrosion rate measured by linear polarimetry in 3.5% by weight aqueous sodium chloride solution is less than about 250% of the rate determined for 15-5 PH (H900 condition). Item 127. An alloy according to item 127.   Less than 1% silicon (Si), less than 0.3% niobium (Nb), less than 0.8% vanadium (V), less than 3% tungsten (W), less than 0.2% titanium (Ti) , Less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), and less than 0.005% boron (B) (percentages by weight) 128. The alloy of claim 1 or claim 127 containing elements.   Less than about 0.02% sulfur (S), less than 0.012% phosphorus (P), less than 0.015% oxygen (O) and less than 0.015% nitrogen (N) (percentages by weight) 128. The alloy of claim 1 or claim 127, comprising:   128. The alloy of claim 1 or claim 127 comprising a substantial lath martensite phase. Cr and Co contained in combination with M _{2 C} carbides to provide a corrosion resistant immobilization layer rich in Cr claim 1 or alloy of claim 127.   128. The alloy of claim 1 or claim 127, further comprising a gettering compound and a grain boundary aggregation strengthening element. La ₂ O ₂ S or _Ce 2 _{O 2} S claim 1 or alloy of claim 127 further comprising a getter ring compound of.   128. The alloy of claim 1 or claim 127, further comprising a grain boundary cohesive strengthening element selected from the group consisting of B, C and Mo. Claim 1 or alloy of claim 127 further comprising an average size of about 10 nanometers less than M _{2 C} carbides precipitate as hydrogen transport inhibitors.   A carbon content of about 10% or less by weight of the alloy is found in the primary MC carbide greater than about 10 nanometers, where M is selected from the group consisting of Ti, V, Nb, Mo, Ta, and combinations thereof 128. The alloy of claim 1 or claim 127. A carbon content of about 2% or less of the alloy by weight is found in carbides greater than about 75 nanometers, and the carbides are M ₆ C, M ₇ C ₃ , M ₂₃ C ₆ , M ₃ C, and M ₂ C. 128. The alloy of claim 1 or claim 127, wherein M is selected from the group consisting of Fe, Cr, Mo, V, W, Nb, Ta, and Ti and combinations thereof.   A carbon content of about 5% or less by weight of the alloy is found in MC carbide greater than about 10 nanometers, and M is from the group consisting of Cr, Mo, V, W, Nb, Ta, Ti and combinations thereof 128. The alloy of claim 1 or claim 127, which is selected.   128. The alloy of claim 1 or 127, wherein the alloy is solution treated at a metal temperature in the range of about 850 <0> C and 1200 <0> C.   128. The alloy of claim 1 or 127, wherein the alloy is solution treated at a metal temperature in the range of about 950 <0> C and 1100 <0> C.   128. The alloy of claim 1 or claim 127, wherein the alloy is cooled from solution treatment to approximately room temperature to form primarily a lath martensite structure.   128. The alloy of claim 1 or claim 127, wherein the alloy is cooled from solution treatment to about room temperature and then further cooled from about room temperature to a metal temperature of less than about -70 ° C to form a primarily lath martensitic structure.   128. The alloy of claim 1 or claim 127, wherein the alloy is cooled from solution treatment to about room temperature and then further cooled from about room temperature to a metal temperature of less than about -195 ° C to form a predominantly lath martensitic structure.   128. The alloy of claim 1 or claim 127, wherein the alloy is tempered in one or more steps at a metal temperature of less than about 600 ° C., and the alloy is cooled between steps to form a primarily lath martensitic structure.   128. The alloy of claim 1 or claim 127, wherein the alloy is tempered in one or more steps at a metal temperature of less than about 300 <0> C and the alloy is cooled between steps to form a primarily lath martensitic structure.   128. The alloy of claim 1 or claim 127, wherein the alloy is tempered in one or more steps at a metal temperature of less than about 400 ° C. and the alloy is cooled between steps to form a primarily lath martensitic structure.   128. The alloy is tempered in one or more steps at a metal temperature in the range of about 400 ° C. and 600 ° C., and the alloy is cooled between steps to form a predominantly lath martensitic structure. Alloy.   128. The alloy is tempered in one or more steps at a metal temperature in the range of about 475 ° C. and 525 ° C., and the alloy is cooled between steps to form a primarily lath martensitic structure. Alloy.   128. The alloy of claim 1 or 127, wherein the alloy is tempered to a hardness greater than about 53 Rockwell C.   128. The alloy of claim 1 or claim 127, wherein the alloy is tempered to a hardness greater than about 50 Rockwell C.   128. The alloy of claim 1 or 127, wherein the alloy is tempered to a hardness greater than about 45 Rockwell C.   128. The alloy of claim 1 or claim 127, wherein the alloy is case hardened to a surface hardness greater than about 67 Rockwell C.   128. The alloy of claim 1 or 127, wherein the alloy is case hardened to a surface hardness of greater than about 60 Rockwell C. Alloy

The toughness / strength ratio, K _Ic / σ _y , greater than or equal to K _Ic / σ _y , where K _Ic is the fracture toughness of the alloy and σ _y is the yield strength. 127 alloy. A method for producing an ultra-high strength corrosion resistant structural steel alloy product comprising:
About 0.1 to 0.3% carbon (C), 8 to 17% cobalt (Co), less than 5% nickel (Ni), greater than 6 and less than 11% chromium (Cr) by weight And combining less than 3% molybdenum (Mo), the balance essentially iron (Fe) and a mixture of molten elements including incidental elements and impurities, and treating the molten mixture to produce an article of manufacture Forming step,
A method for producing an ultra-high-strength, corrosion-resistant structural steel alloy product.   The steel alloy product comprises less than about 1% silicon (Si), less than 0.3% nickel (Nb), less than 0.8% vanadium (V), less than 3% tungsten (W), 0.2% Less than% titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), and less than 0.005% boron (B) (percentages are by weight 60. The method according to claim 59, wherein said method is formulated to contain one or more elements selected from the group comprising:   The steel alloy product comprises less than about 0.02% sulfur (S), less than 0.012% phosphorus (P), less than 0.015% oxygen (O) and less than 0.015% nitrogen (N). 60. The method according to claim 59, formulated to contain (percentages by weight). Processing the steel alloy product,
(A) homogenization of the steel alloy article;
(B) hot working of the steel alloy article;
60. The method according to claim 59, comprising: (c) normalizing the steel alloy article; and (d) annealing the steel alloy article.   63. The method according to claim 62, wherein the homogenization is at least 4 hours at a metal temperature in the range of about 1100 <0> C to 1400 <0> C.   63. The method according to claim 62, wherein the homogenization is at least 4 hours at a metal temperature in the range of about 1200 <0> C to 1300 <0> C.   63. The method according to claim 62, wherein the hot working is at a metal temperature in the range of about 840 ° C to 1300 ° C and results in a total reduction in cross-sectional area of at least about 5 to 1.   63. The method according to claim 62, wherein the hot working is at a metal temperature in the range of about 1030 ° C to 1200 ° C and results in a total reduction in cross-sectional area of at least about 5 to 1.   63. The method according to claim 62, wherein the normalizing is at a metal temperature in the range of about 880 ° C to 1100 ° C.   63. The method according to claim 62, wherein the normalization is at a metal temperature in the range of about 980 ° C to 1080 ° C.   63. The method according to claim 62, wherein the annealing is longer than 1 hour at a metal temperature in the range of about 600 <0> C to 850 <0> C.   63. The method according to claim 62, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 650 <0> C to 790 <0> C. Processing the steel alloy product,
(A) homogenization of the steel alloy article;
60. The method according to claim 59, comprising: (b) hot working of the steel alloy article; and (c) annealing the steel alloy article.   72. The method according to claim 71, wherein said homogenization is at least 4 hours at a metal temperature in the range of about 1100 <0> C to 1400 <0> C.   72. The method according to claim 71, wherein said homogenization is at least 4 hours at a metal temperature in the range of about 1200 <0> C to 1300 <0> C.   72. The method according to claim 71, wherein the hot working is at a metal temperature in the range of about 840 <0> C to 1300 <0> C and produces a total reduction in cross-sectional area of at least about 5 to 1.   72. The method according to claim 71, wherein said hot working is at a metal temperature in the range of about 1030 <0> C to 1200 <0> C and results in a total reduction in cross-sectional area of at least about 5 to 1.   72. The method according to claim 71, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 600 <0> C to 850 <0> C.   72. The method according to claim 71, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 650 <0> C to 790 <0> C. The steel alloy article,
(A) Solution treatment of the steel alloy article,
(B) cooling the steel alloy article; and (c) tempering the steel alloy article;
63. The method according to claim 62, further processed by the steps of:   The method according to claim 78, wherein said solution treatment is at a metal temperature in the range of about 850 ° C to 1100 ° C.   The method according to claim 78, wherein said solution treatment is at a metal temperature in the range of about 950 ° C to 1050 ° C.   79. A method according to claim 78, wherein the cooling is to approximately room temperature.   The method according to claim 78, wherein said cooling is to a metal temperature of less than about -70 ° C.   The method according to claim 78, wherein said cooling is to a metal temperature of less than about -195 ° C.   79. The method according to claim 78, wherein the tempering is at a metal temperature of less than about 600 ° C. in one or more steps and the steel alloy product is cooled between steps.   79. The method according to claim 78, wherein the tempering is at a metal temperature of less than about 500 ° C. in one or more steps and the steel alloy product is cooled between steps.   79. The method according to claim 78, wherein the tempering is at a metal temperature of less than about 400 ° C. in one or more steps and the steel alloy product is cooled between steps.   79. The method according to claim 78, wherein the tempering is at a metal temperature of less than about 300 <0> C in one or more steps and the steel alloy product is cooled between steps.   79. The method according to claim 78, wherein said tempering is at a metal temperature in the range of about 400 <0> C to 600 <0> C in one or more steps and cools the steel alloy product between steps.   79. The method according to claim 78, wherein the tempering is at a metal temperature in the range of about 450 ° C. to 540 ° C. in one or more steps and the steel alloy product is cooled between steps. The steel alloy article,
(A) Solution treatment of the steel alloy article,
(B) cooling the steel alloy article; and (c) tempering the steel alloy article;
72. The method according to claim 71, further processed by:   The method according to claim 90, wherein said solution treatment is at a metal temperature in the range of about 850 ° C to 1100 ° C.   The method according to claim 90, wherein said solution treatment is at a metal temperature in the range of about 950 ° C to 1050 ° C.   The method according to claim 90, wherein said cooling is to a metal temperature of about room temperature.   The method according to claim 90, wherein said cooling is to a metal temperature of less than about -70 ° C.   The method according to claim 90, wherein said cooling is to a metal temperature of less than about -195 ° C.   95. The method according to claim 90, wherein the tempering is at a metal temperature of less than about 600 ° C. in one or more steps and the steel alloy product is cooled between steps.   The method according to claim 90, wherein said tempering is at a metal temperature of less than about 500 ° C in one or more steps and cooling the steel alloy product between steps.   The method according to claim 90, wherein said tempering is at a metal temperature of less than about 400 ° C in one or more steps and cooling the steel alloy product between steps.   94. The method according to claim 90, wherein the tempering is at a metal temperature of less than about 300 <0> C in one or more steps and the steel alloy product is cooled between steps.   The method according to claim 90, wherein said tempering is at a metal temperature in the range of about 400 ° C to 600 ° C in one or more steps and cooling the steel alloy product between steps.   The method according to claim 90, wherein said tempering is at a metal temperature in the range of about 450 ° C to 540 ° C in one or more steps and cooling the steel alloy product between steps. Process includes the step of forming a major _{M 2} C carbide in the alloy, wherein, M is Cr, Mo, V, W, Nb, is selected from the group consisting of Ta and combinations thereof, The method according to claim 59 .   60. The method according to claim 59, wherein the treatment comprises a heat treatment to form a substantially martensitic phase material. The treatment includes a heat treatment that forms a majority of carbon by weight as M ₂ C carbide, wherein M is from the group consisting of Cr, Fe, Mo, V, W, Nb, Ta, Ti, and combinations thereof 60. The method according to claim 59, wherein the method is selected.   About 0.2 to 0.26% carbon (C) by weight, 11 to 15% cobalt (Co), 2.0 to 3.0% nickel (Ni), 7.5 to 9.5% by weight Combining chromium (Cr), 1.0 to 2.0% molybdenum (Mo), and less than 0.8% vanadium (V), the balance being substantially iron (Fe) and incidental elements and impurities Alloy composition to contain.   About 0.20 to 0.25% by weight of carbon (C), 12 to 15% cobalt (Co), 2.0 to 3.0% nickel (Ni), 7.0 to 9.0% Chromium (Cr), 1.0 to 3.0% molybdenum (Mo), less than 2.5% tungsten (W), less than 0.75% silicon (Si), and less than 0.8% vanadium ( V), an alloy composition containing a balance of essentially iron (Fe) and a combination of incidental elements and impurities.   About 0.10 to 0.20% by weight of carbon (C), 12 to 17% cobalt (Co), 2.5 to 5.0% nickel (Ni), 8.5 to 9.5% Combining chromium (Cr), 1.0 to 2.0% molybdenum (Mo), and less than 0.8% vanadium (V), the balance being essentially iron (Fe) and incidental elements and impurities Alloy composition to contain.   About 0.25 to 0.28% carbon (C) by weight, 11 to 15% cobalt (Co), 1.0 to 3.0% nickel (Ni), 7.0 to 9.0% by weight Chromium (Cr), less than 1.0% molybdenum (Mo), less than 1.0% silicon (Si), and less than 0.8% vanadium (V), the balance being essentially iron (Fe) and An alloy composition containing a combination of incidental elements and impurities.   About 0.22 to 0.25% by weight carbon (C), 12 to 13% cobalt (Co), 2.5 to 3.0% nickel (Ni), 8.5 to 9.5% Combining chromium (Cr), 1.0 to 1.5% molybdenum (Mo), and less than 0.8% vanadium (V), the balance being essentially iron (Fe) and incidental elements and impurities Alloy composition to contain.   About 0.18 to 0.25% by weight of carbon (C), 10 to 15% cobalt (Co), 4.0 to 8.0% nickel (Ni), 8.0 to 10.0% Chromium (Cr), 1.0 to 3.0% molybdenum (Mo), less than 3.0% tungsten (W), less than 1.0% silicon (Si), and less than 0.8% vanadium ( V), an alloy composition containing a balance of essentially iron (Fe) and a combination of incidental elements and impurities.   About 0.18 to 0.25% carbon (C) by weight, 6 to 10% cobalt (Co), 4.0 to 8.0% nickel (Ni), 8.0 to 10.0% by weight Chromium (Cr), 1.0 to 3.0% molybdenum (Mo), less than 3.0% tungsten (W), less than 1.0% silicon (Si), and less than 0.8% vanadium ( V), an alloy composition containing a balance of essentially iron (Fe) and a combination of incidental elements and impurities.   About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), 0 to 5% nickel (Ni), 6 to 12% chromium (Cr), less than 1% Silicon (Si), less than 0.5% manganese (Mn), and less than 0.15% copper (Cu), approximately less than 3% molybdenum (Mo), less than 0.3% niobium (Nb) An additive selected from the group comprising less than 0.8% vanadium (V), less than 0.2% tantalum (Ta), less than 3% tungsten (W), and combinations thereof, approximately 0.2 % Less than titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), less than 0.005% boron (B), and combinations thereof Additional additives selected from the group comprising, and essentially (Fe) is a balance and the alloy composition containing in combination with incidental elements and impurities.   About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), 0 to 5% nickel (Ni), 6 to 12% chromium (Cr), less than 1% Silicon (Si), less than 0.5% manganese (Mn), and less than 0.15% copper (Cu), approximately less than 3% molybdenum (Mo), less than 0.3% niobium (Nb) An additive selected from the group comprising less than 0.8% vanadium (V), less than 0.2% tantalum (Ta), less than 3% tungsten (W), and combinations thereof, approximately 0.2 % Less than titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), less than 0.005% boron (B), and combinations thereof About 0.02% additional additive selected from the group comprising Sulfur (S), 0.012% phosphorus (P), 0.015% oxygen (O) and 0.015% nitrogen (N) impurities, the balance being essentially iron (Fe) and incidental An alloy composition containing a combination of an element and an impurity.   About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), 0 to 10% nickel (Ni), 6 to 12% chromium (Cr), less than 1% Silicon (Si), less than 0.5% manganese (Mn), and less than 0.15% copper (Cu), approximately less than 3% molybdenum (Mo), less than 0.3% niobium (Nb) An additive selected from the group comprising less than 0.8% vanadium (V), less than 0.2% tantalum (Ta), less than 3% tungsten (W), and combinations thereof, approximately 0.2 % Less than titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), less than 0.005% boron (B), and combinations thereof Additional additives selected from the group comprising, and essentially Iron (Fe) is a balance and the alloy composition containing in combination with incidental elements and impurities.   About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), 0 to 10% nickel (Ni), 6 to 12% chromium (Cr), less than 1% Silicon (Si), less than 0.5% manganese (Mn), and less than 0.15% copper (Cu), approximately less than 3% molybdenum (Mo), less than 0.3% niobium (Nb) An additive selected from the group comprising less than 0.8% vanadium (V), less than 0.2% tantalum (Ta), less than 3% tungsten (W), and combinations thereof, approximately 0.2 % Less than titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), less than 0.005% boron (B), and combinations thereof About 0.02% additional additive selected from the group comprising Full sulfur (S), 0.012% phosphorus (P), 0.015% oxygen (O) and 0.015% nitrogen (N) impurities, the balance being essentially iron (Fe) and Alloy composition containing in combination with incidental elements and impurities. Alloys containing M ₂ C carbide with a diameter less than about 10 nanometers have a carbon content greater than about 85% by weight, where M is Cr, Mo, V, W, Nb, Ta, and combinations thereof 116. The alloy according to any of claims 105 to 115, selected from the group consisting of: Alloys containing M ₂ C carbide with a diameter less than about 5 nanometers have a carbon content greater than about 85% by weight, where M is Cr, Mo, V, W, Nb, Ta and combinations thereof 116. The alloy according to any of claims 105 to 115, selected from the group consisting of:   116. The alloy of any of claims 105 to 115, having a maximum tensile strength greater than about 240 ksi.   116. The alloy of any of claims 105 to 115, having a yield strength greater than about 200 ksi.   116. An alloy according to any of claims 105 to 115 comprising dispersed metal (M) carbide particles, wherein the particles have the formula MXC, where X ≦ 2 for most of the weight percent of the particles. An alloy that is mainly in the martensitic phase.   The alloy is in a martensitic phase and includes dispersed metal carbide, the metal carbide having a nominal dimension less than about 10 nanometers in diameter, and mainly in the range of about 2 to 1 or less metal ions to carbon. 116. An alloy according to any of claims 105 to 115, having an ionic ratio.   The alloy is in the martensitic phase and includes dispersed metal carbide, the metal carbide having a nominal dimension less than about 5 nanometers in diameter, and mainly in the range of about 2 to 1 or less metal ions to carbon. 116. An alloy according to any of claims 105 to 115, having an ionic ratio.   The alloy has dispersed metal carbide, wherein the ratio of metal ions to carbon ions is primarily about 2 to 1, and the metals are Cr, Mo, V, W, Nb, Ta, Ti, and their 116. The alloy according to any of claims 105 to 115, selected from the group consisting of combinations.   The alloy has dispersed metal carbide, the metal is selected from the group consisting of Cr, Mo, V, W, Nb, Ta, Ti, and the ratio of metal ions to carbon ions is primarily about 2 to 1, 116. An alloy according to any of claims 105 to 115, wherein the alloy is substantially in the martensitic phase.   116. The alloy of any of claims 105 to 115, wherein the alloy has a nominal particle size that is approximately equal to or smaller than ASTM particle size number 5 (ASTM E112).   116. The alloy of any of claims 105-115, wherein the alloy is primarily in the martensitic phase and has a nominal particle size that is approximately equal to or smaller than ASTM particle size No. 5 (ASTM E112).   About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), less than 10% nickel (Ni), greater than 6 and less than 11% chromium (Cr); And less than 3% molybdenum (Mo), the balance being essentially iron (Fe), and a combination of incidental elements and impurities. A method for producing an ultra-high-strength, corrosion-resistant structural steel alloy product,
About 0.1 to 0.3% carbon (C) by weight, 8 to 17% cobalt (Co), less than 10% nickel (Ni), greater than 6 and less than 11% chromium (Cr); And combining less than 3% molybdenum (Mo), the balance being essentially iron (Fe), and a mixture of molten elements containing incidental elements and impurities, and treating the molten mixture to produce an article of manufacture A method comprising the step of forming.   The steel alloy product is approximately less than 1% silicon (Si), less than 0.3% niobium (Nb), less than 0.8% vanadium (V), less than 3% tungsten (W), 0.2 Less than% titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements, less than 0.15% zirconium (Zr), and less than 0.005% boron (B) (percentages are by weight 129. The method according to claim 128, wherein said compound is formulated to contain one or more elements from the group comprising:   The steel alloy product is approximately less than 0.02% sulfur (S), less than 0.012% phosphorus (P), less than 0.015% oxygen (O) and less than 0.015% nitrogen (N). 129. A method according to claim 128, formulated to contain (percentages by weight). Processing the steel alloy product,
(A) homogenization of the steel alloy article,
(B) hot working of the steel alloy article;
(C) normalizing the steel alloy article; and (d) annealing the steel alloy article;
129. A method according to claim 128, comprising:   132. The method according to claim 131, wherein the homogenization is at least 4 hours at a metal temperature in the range of about 1100 <0> C to 1400 <0> C.   132. The method according to claim 131, wherein said homogenization is at least 4 hours at a metal temperature in the range of about 1200 <0> C to 1300 <0> C.   132. The method according to claim 131, wherein said hot working is at a metal temperature in the range of about 840 <0> C to 1300 <0> C and results in a total reduction in cross-sectional area of at least about 5 to 1.   132. The method according to claim 131, wherein said hot working is at a metal temperature in the range of about 1030 <0> C to 1200 <0> C and results in a total reduction in cross-sectional area of at least about 5 to 1.   132. The method according to claim 131, wherein said normalizing is at a metal temperature in the range of about 880 ° C to 1100 ° C.   132. The method according to claim 131, wherein the normalization is at a metal temperature in the range of about 980 ° C to 1080 ° C.   132. The method according to claim 131, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 600 <0> C to 850 <0> C.   132. The method according to claim 131, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 650 <0> C to 790 <0> C. Processing the steel alloy product,
(A) homogenization of the steel alloy article;
129. The method according to claim 128, comprising: (b) hot working of the steel alloy article; and (c) annealing the steel alloy article.   141. The method according to claim 140, wherein the homogenization is at least 4 hours at a metal temperature in the range of about 1100 <0> C to 1400 <0> C.   141. The method according to claim 140, wherein the homogenization is at least 4 hours at a metal temperature in the range of about 1200 <0> C to 1300 <0> C.   141. The method according to claim 140, wherein the hot working is at a metal temperature in the range of about 840 ° C to 1300 ° C and results in a total reduction in cross-sectional area of at least about 5 to 1.   141. The method according to claim 140, wherein the hot working is at a metal temperature in the range of about 1030 ° C to 1200 ° C and results in a total reduction in cross-sectional area of at least about 5 to 1.   141. The method according to claim 140, wherein the annealing is longer than 1 hour at a metal temperature in the range of about 600 <0> C to 850 <0> C.   141. The method according to claim 140, wherein said annealing is longer than 1 hour at a metal temperature in the range of about 650 <0> C to 790 <0> C. The steel alloy article,
(A) Solution treatment of the steel alloy article,
132. The method according to claim 131, further processed by the steps of (b) cooling the steel alloy article, and (c) tempering the steel alloy article.   148. The method according to claim 147, wherein the solution treatment is at a metal temperature in the range of about 850 ° C to 1100 ° C.   148. The method according to claim 147, wherein the solution treatment is at a metal temperature in the range of about 950 ° C to 1050 ° C.   148. The method according to claim 147, wherein the cooling is to approximately room temperature.   148. The method according to claim 147, wherein the cooling is at a metal temperature less than about -70C.   148. The method according to claim 147, wherein the cooling is at a metal temperature less than about -195C.   148. The method according to claim 147, wherein the tempering is at a metal temperature of less than about 600 ° C. in one or more steps and the steel alloy product is cooled between steps.   148. The method according to claim 147, wherein the tempering is at a metal temperature of less than about 500 ° C. in one or more steps and the steel alloy product is cooled between steps.   148. The method according to claim 147, wherein the tempering is at a metal temperature of less than about 400 ° C. in one or more steps and the steel alloy product is cooled between steps.   148. The method according to claim 147, wherein the tempering is at a metal temperature of less than about 300 <0> C in one or more steps and the steel alloy product is cooled between steps.   148. The method according to claim 147, wherein the tempering is at a metal temperature in the range of about 400 ° C. to 600 ° C. in one or more steps and the steel alloy product is cooled between steps.   148. The method according to claim 147, wherein the tempering is at a metal temperature in the range of about 450 ° C. to 540 ° C. in one or more steps and the steel alloy product is cooled between steps. The steel alloy article,
(A) Solution treatment of the steel alloy article,
(B) cooling the steel alloy article; and (c) tempering the steel alloy article;
145. The method according to claim 140, further processed by:   160. The method according to claim 159, wherein the solution treatment is at a metal temperature in the range of about 850 <0> C to 1100 <0> C.   160. The method according to claim 159, wherein the solution treatment is at a metal temperature in the range of about 950 <0> C to 1050 <0> C.   160. The method according to claim 159, wherein the cooling is to a metal temperature of about room temperature.   The method according to claim 159, wherein said cooling is to a metal temperature of less than about -70 ° C.   The method according to claim 159, wherein said cooling is to a metal temperature of less than about -195 ° C.   159. The method according to claim 159, wherein said tempering is at a metal temperature of less than about 600 ° C. in one or more steps and cooling the steel alloy product between steps.   159. The method according to claim 159, wherein the tempering is at a metal temperature of less than about 500 ° C. in one or more steps and the steel alloy product is cooled between steps.   160. The method according to claim 159, wherein said tempering is at a metal temperature of less than about 400 ° C. in one or more steps and cooling the steel alloy product between steps.   160. The method according to claim 159, wherein the tempering is at a metal temperature of less than about 300 <0> C in one or more steps and the steel alloy product is cooled between steps.   160. The method according to claim 159, wherein said tempering is at a metal temperature in the range of about 400 <0> C to 600 <0> C in one or more steps and cools the steel alloy product between steps.   160. The method according to claim 159, wherein said tempering is at a metal temperature in the range of about 450 [deg.] C to 540 [deg.] C in one or more steps and cools the steel alloy product between steps. Process includes the step of forming a major _{M 2} C carbide in the alloy, wherein, M is Cr, Mo, V, W, Nb, is selected from the group consisting of Ta and combinations thereof, The method according to claim 128 .   129. The method according to claim 128, wherein the treatment comprises a heat treatment that forms a substantially martensitic phase material. The treatment includes a heat treatment that forms a majority of carbon by weight as M ₂ C carbide, wherein M is from the group consisting of Cr, Fe, Mo, V, W, Nb, Ta, Ti, and combinations thereof 129. The method according to claim 128, which is selected.

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