# Patent

## Title
Classes of modal structured steel with static refinement and dynamic strengthening and method of making thereof

## Abstract

The present disclosure is directed at formulations and methods to provide new steel alloys having relatively high strength and ductility. The alloys may be provided in sheet or pressed form and characterized by their particular alloy chemistries and identifiable crystalline grain size morphology. The alloys are such that they include boride grains present as pinning phases. Mechanical properties of the alloys in what is termed a Class 1 Steel indicate yield strengths of 300 MPa to 840 MPa, tensile strengths of 630 to 1100 MPa and elongations of 10% to 40%. In what is termed a Class 2 steel, the alloys indicate yield strengths of 300 MPa to 1300 MPa, tensile strengths of 720 MPa to 1580 MPa and elongations of 5% to 35%.

## Background
CROSS REFERENCE TO RELATED APPLICATION
This application claims the benefit of U.S. Provisional Application Ser. No. 61/488,558 filed May 20, 2011 and U.S. Provisional Application Ser. No. 61/586,951 filed Jan. 16, 2012, the teachings of which are incorporated herein by reference.


FIELD OF INVENTION
This application deals with new modal structured steel alloys with application to a sheet production by chill surface processing. Two new classes of steel are provided involving the achievement of various levels of strength and ductility. Three new structure types have been identified which may be achieved by disclosed mechanisms.
BACKGROUND
Steels have been used by mankind for at least 3,000 years and are widely utilized in industry comprising over 80% by weight of all metallic alloys in industrial use. Existing steel technology is based on manipulating the eutectoid transformation. The first step is to heat up the alloy into the single phase region (austenite) and then cool or quench the steel at various cooling rates to form multiphase structures which are often combinations of ferrite, austenite, and cementite. Depending on how the steel is cooled, a wide variety of characteristic microstructures (i.e. pearlite, bainite, and martensite) can be obtained with a wide range of properties. This manipulation of the eutectoid transformation has resulted in the wide variety of steels available nowadays.
Currently, there are over 25,000 worldwide equivalents in 51 different ferrous alloy metal groups. For steel, which is produced in sheet form, broad classifications may be employed based on tensile strength characteristics. Low Strength Steels (LSS) may be defined as exhibiting tensile strengths less than 270 MPa and include types such as interstitial free and mild steels. High-Strength Steels (HSS) may be steel defined as exhibiting tensile strengths from 270 to 700 MPa and include types such as high strength low alloy, high strength interstitial free and bake hardenable steels. Advanced High-Strength Steels (AHSS) steels may have tensile strengths greater than 700 MPa and include types such as martensitic steels (MS), dual phase (DP) steels, transformation induced plasticity (TRIP) steels, and complex phase (CP) steels. As the strength level increases, the ductility of the steel generally decreases. For example, LSS, HSS and AHSS may indicate tensile elongations at levels of 25%-55%, 10%-45% and 4%-30%, respectively.
SUMMARY
The present disclosure relates to a method for producing a metallic alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.50 atomic percent, B at 4.00 to 8.00 atomic percent, Si at 4.00 to 8.00 atomic percent. This may then be followed by melting the alloy and solidifying to provide a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm. On may then mechanically stress the alloy and/or heat to form at least one of the following grain size distributions and mechanical property profiles, wherein the boride grains provide pinning phases that resist coarsening of said matrix grains:
(a) matrix grain size in the range of 500 nm to 20,000 nm, boride grain size in the range of 25 nm to 500 nm, precipitation grain size in the range of 1 nm to 200 nm wherein the alloy indicates a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa and tensile elongation of 10 to 40%; or
(b) matrix grain size in the range of 100 nm to 2000 nm and boride grain size in the range of 25 nm to 500 nm which has a yield strength of 300 MPa to 600 MPa.
The present disclosure also relates to a method for producing a metallic alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.5 atomic percent, B at 4.0 to 8.0 atomic percent, Si at 4.0 to 8.0 atomic percent. This may be followed by melting the alloy and solidifying to provide a matrix grain size of 500 nm to 20,000 nm containing 10% to 70% by volume ferrite and a boride grain size in the range of 25 nm to 500 nm wherein the boride grains provide pinning phases that resist coarsening of the matrix grains upon application of heat and wherein the alloy has a yield strength of 300 MPa to 600 MPa. This may then be followed by heating the alloy wherein the grain size is in the range of 100 nm to 2000 nm, the boride grain size remains in the range of 25 nm to 500 nm and the level of ferrite increases to 20% to 80% by volume. One may then stress the alloy to a level that exceeds the yield strength of 300 MPa to 600 MPa wherein the grain size remains in the range of 100 nm to 2000 nm, the boride grain size remains in the range of 25 nm to 500 nm, along with the formation of precipitation grains in the range of 1 nm to 200 nm and the alloy has a tensile strength of 720 MPa to 1580 MPa and an elongation of 5% to 35%.
The present disclosure also relates to a metallic alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.5 atomic percent, B at 4.0 to 8.0 atomic percent, and Si at 4.0 to 8.0 atomic percent. The alloy indicates a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm wherein the alloy indicates at least one of the following:
(a) upon exposure to mechanical stress the alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa, and tensile elongation of 10 to 40%; or
(b) upon exposure to heat, followed by mechanical stress, the alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 1300 MPa, tensile strength of 720 MPa to 1580 MPa, and tensile elongation of 5.0% to 35.0%.
The present disclosure also relates to a metallic alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.5 atomic percent, B at 4.0 to 8.0 atomic percent and Si at 4.0 to 8.0 atomic percent. The alloy indicates a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm wherein the alloy indicates at least one of the following:
(a) upon exposure to mechanical stress the alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa, tensile elongation of 10% to 40%, a matrix grain size in the range of 500 nm to 20,000 nm, a boride grain size in the range of 25 nm to 500 nm and a precipitation grain size in the range of 1.0 nm to 200 nm; or
(b) upon exposure to heat followed by mechanical stress, the alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 1300 MPa, tensile strength of 720 MPa to 1580 MPa, tensile elongation of 5% to 35% and a matrix grain size in the range of 100 nm to 2000 nm, a boride grain size in the range of 25 nm to 500 nm, and a precipitation grain size in the range of 1 nm to 200 nm.



BRIEF DESCRIPTION OF THE DRAWINGS
The detailed description below may be better understood with reference to the accompanying FIGS. which are provided for illustrative purposes and are not to be considered as limiting any aspect of this invention.
FIG. 1 illustrates an exemplary twin-roll process.
FIG. 2 illustrates an exemplary thin slab casting process.
FIG. 3A illustrates structures and mechanisms regarding the formation of Class 1 Steel herein.
FIG. 3B illustrates structures and mechanism regarding the formation of Class 2 Steel herein.
FIG. 3C illustrates a general scheme for the formation of Class 1 and Class 2 Steel herein.
FIG. 4 illustrates a representative stress-strain curve of material containing modal phase formation.
FIG. 5 illustrates a representative stress-strain curve for the indicated structures and associated mechanisms of formation.
FIG. 6 illustrates a photograph of Alloy 19 sheet under specified conditions.
FIG. 7 illustrates a comparison of stress-strain curves of indicated steel types as compared to Dual Phase (DP) steels.
FIG. 8 illustrates a comparison of stress-strain curves of indicated steel types as compared to Complex Phase (CP) steels.
FIG. 9 illustrates a comparison of stress-strain curves of indicated steel types as compared to Transformation Induced Plasticity (TRIP) steels.
FIG. 10 illustrates a comparison of stress-strain curves of indicated steel-types as compared to Martensitic (MS) steels.
FIG. 11 illustrates a SEM micrograph of Modal Structure herein of Alloy 2.
FIG. 12 illustrates a SEM micrograph of Modal Structure herein of Alloy 11 after HIP cycle at 1000° C. for 1 hour.
FIG. 13 illustrates a SEM micrograph of Modal Structure herein of Alloy 18 after HIP cycle at 1100° C. for 1 hour.
FIG. 14 illustrates a SEM micrograph of Modal Structure of Alloy 1 after HIP cycle at 1000° C. for 1 hour and annealing at 350° C. for 20 minutes.
FIG. 15 is an SEM micrograph of Modal Structure herein in Alloy 14.
FIG. 16 is picture of as-cast Alloy 1 sheet.
FIG. 17 is an SEM backscattered electron micrograph of Alloy 1 under the indicated conditions of formation.
FIG. 18 is X-ray diffraction data for Alloy 1 sheet.
FIG. 19 is X-ray diffraction data for Alloy 1 sheet in the HIPed condition.
FIG. 20 is X-ray diffraction data for Alloy 1 sheet in the HIPed condition.
FIG. 21 is TEM micrographs of Alloy 1 under the indicated conditions.
FIG. 22 is a stress-strain plot of Alloy 1 under the indicated conditions of formation.
FIG. 23 is a comparison of X-ray data for Alloy 1 under the indicated conditions.
FIG. 24 is X-ray diffraction data for the gage section of tensile tested sample from Alloy 1 in the HIPed condition.
FIG. 25 is a calculated X-ray diffraction pattern for iron based hexagonal phase in the gage section of tensile tested sample from Alloy 1 sheet.
FIG. 26 is a TEM micrograph of Alloy 1 sheet HIPed under the indicated conditions.
FIG. 27 is a TEM micrograph of the gage section microstructure in a tensile tested specimen from Alloy 1 sheet under the indicated conditions.
FIG. 28 is a TEM micrograph of the gage section microstructure in tensile tested specimen from Alloy 1 sheet under the indicated conditions.
FIG. 29 is a picture of as-cast Alloy 14 sheet.
FIG. 30 is a SEM backscattered electron micrograph of the Alloy 14 sheet under the indicated conditions.
FIG. 31 X-ray diffraction data for Alloy 14 sheet under the indicated conditions.
FIG. 32 is X-ray diffraction data for Alloy 14 in the HIPed condition.
FIG. 33 is X-ray diffraction data for Alloy 14 in the HIPed condition.
FIG. 34 are TEM micrographs of the Alloy 14 sheet under the indicated conditions.
FIG. 35 is a stress-strain plot of Alloy 14 sheet under the indicated conditions.
FIG. 36 is a comparison of X-ray data for Alloy 14 sheet under the indicated conditions.
FIG. 37 is X-ray diffraction data from the gage section of tensile tested sample from Alloy 14 in the HIPed condition.
FIG. 38 is a calculated X-ray diffraction pattern for iron based hexagonal phase in the gage section of tensile tested sample from Alloy 14 sheet in the HIPed condition.
FIG. 39 is a TEM micrograph of Alloy 14 sheet HIPed at 1000° C. under the indicated conditions.
FIG. 40 is a TEM micrograph of the Alloy 14 tensile tested gage specimen under the indicated conditions.
FIG. 41 is a picture of as-case Alloy 19 sheet.
FIG. 42 is a SEM backscattered electron micrograph of Alloy 19 sheet under the indicated conditions.
FIG. 43 is X-ray diffraction data for Alloy 19 sheet under the indicated conditions.
FIG. 44 is X-ray diffraction data for Alloy 19 sheet in the HIPed condition.
FIG. 45 is X-ray diffraction data for Alloy 19 sheet in the HIPed condition.
FIG. 46 is TEM electron micrographs of the Alloy 19 sheet under the indicated conditions.
FIG. 47 is a stress-strain plot of Alloy 19 sheet under the indicated conditions.
FIG. 48 is a comparison between X-ray data for Alloy 19 sheet after the HIP cycle at 1100° C. for 1 hour and heat treatment at 700° C. for 20 minutes.
FIG. 49 is X-ray diffraction data for the gage section of tensile tested sample from Alloy 19 under the indicated conditions.
FIG. 50 is a calculated X-ray diffraction pattern for iron based hexagonal phase found in the gage section of tensile tested sample from Alloy 19 under the indicated conditions.
FIG. 51 is a TEM micrograph of Alloy 19 under the indicated conditions.
FIG. 52 is a TEM micrograph of Alloy 19 tensile tested gage specimen under the indicated conditions.
FIG. 53 is a TEM micrograph of Alloy 19 tensile tested gage specimen under the indicated conditions.
FIG. 54(a) illustrates stain hardening in alloy sheets with different mechanisms of structural formation.
FIG. 54(b) illustrates tensile properties for the sheets in FIG. 54(a).
FIG. 55 is a stress-strain curve for Alloy 1 sheet at different strain rates.
FIG. 56 is a stress-strain curve for Alloy 19 at different strain rates.
FIG. 57 is a stress-strain curve for Alloy 19 sheet under the indicated conditions.
FIG. 58(a) is a stress-strain curve for Alloy 19 sheet after prestraining to 10%.
FIG. 58(b) is a stress-strain curve for Alloy 19 sheet after prestraining to 10% and subsequent annealing at 1150° C. for 1 hour.
FIG. 59 is a stress-strain curve for Alloy 19 under the indicated conditions.
FIG. 60 illustrates the sample geometry of Alloy 19 under the indicated conditions.
FIG. 61 is a SEM image of microstructure of the gage section of the tensile specimens of Alloy 19 under the indicated conditions.
FIG. 62 is a SEM image of the gage section of the tensile specimens from Alloy 19 under the indicated conditions.
FIG. 63(a) is a plane view of the plate of Alloy 3 after Erichsen test stopped at maximum load.
FIG. 63(b) is a side view of the plate of Alloy 3 after Erichsen test stopped at maximum load.
FIG. 64 is a photograph of the as-cast sheets from Alloy 1 at three different thicknesses.
FIG. 65 is an example of a stress-strain curve of the indicated selected alloys.
FIG. 66 is a stress-strain curve of ductile melt-spun ribbon of test Alloy 47.



DETAILED DESCRIPTION
Steel Strip/Sheet Sizes
Through chill surface processing, steel sheet, as described in this application, with thickness in range of 0.3 mm to 150 mm can be produced in cast thickness and with widths in the range of 100 to 5000 mm. These thickness ranges and width ranges may be adjusted in these ranges to 0.1 mm increments. Preferably, one may use twin roll casting which can provide sheet production at thicknesses from 0.3 to 5 mm and from 100 mm to 5000 mm in width. Preferably, one may also utilize thin slab casting which can provide sheet production at thicknesses from 0.5 to 150 mm and from 100 mm to 5000 mm in width. Cooling rates in the sheet would be dependent on the process but may vary from 11×103 to 4×10−2K/s. Cast parts through various chill surface methods with thickness up to 150 mm, or in the range of 1 mm to 150 mm are also contemplated herein from various methods including, permanent mold casting, investment casting, die casting, etc. Also, powder metallurgy through either conventional press and sintering or through HIPing/forging is a contemplated route to make partial or fully dense parts and devices utilizing the chemistries, structures, and mechanisms described in this application (i.e. the Class 1 or Class 2 Steel described herein).
Production Routes
Twin Roll Casting Description
One of the examples of steel production by chill surface processing would be the twin roll process to produce steel sheet. A schematic of the Nucor/Castrip process is shown in FIG. 1. As shown, the process can be broken up into three stages; Stage 1—Casting, Stage 2—Hot Rolling, and Stage 3—Strip Coiling. During Stage 1, the sheet is formed as the solidifying metal is brought together in the roll nip between the rollers which are generally made out of copper or a copper alloy. Typical thickness of the steel at this stage is 1.7 to 1.8 mm in thickness but by changing the roll separation distance can be varied from 0.8 to 3.0 mm in thickness. During Stage 2, the as-produced sheet is hot rolled, typically from 700 to 1200° C. which acts to eliminate macrodefects such as the formation of pores, dispersed shrinkage, blowholes, pinholes, slag inclusions etc. from the production process as well as allowing solutionizing of key alloying elements, austenitization, etc. The thickness of the hot rolled sheet can be varied depending on the targeted market but is generally in the range from 0.3 to 2.0 mm in thickness. During Stage 3, the temperature of the sheet and time at temperature typically from 300 to 700° C. can be controlled by adding water cooling and changing the length of the run-out of the sheet prior to coiling. Besides hot rolling, Stage 2 could also be done by alternate thermomechanical processing strategies such as hot isostatic processing, forging, sintering etc. Stage 3, besides controlling the thermal conditions during the strip coiling process, could also be done by post processing heat treating in order to control the final microstructure in the sheet.
Thin Slab Casting Description
Another example of steel production by chill surface processing would be the thin slab casting process to produce steel sheet. A schematic of the Arvedi ESP process is shown in FIG. 2. In an analogous fashion to the twin roll process, the thin slab casting process can be separated into three stages. In Stage 1, the liquid steel is both cast and rolled in an almost simultaneous fashion. The solidification process begins by forcing the liquid melt through a copper or copper alloy mold to produce initial thickness typically from 50 to 110 mm in thickness but this can be varied (i.e. 20 to 150 mm) based on liquid metal processability and production speed. Almost immediately after leaving the mold and while the inner core of the steel sheet is still liquid, the sheet undergoes reduction using a multistep rolling stand which reduces the thickness significantly down to 10 mm depending on final sheet thickness targets. In Stage 2, the steel sheet is heated by going through one or two induction furnaces and during this stage the temperature profile and the metallurgical structure is homogenized. In Stage 3, the sheet is further rolled to the final gage thickness target which may be in the 0.5 to 15 mm thickness range. Immediately after rolling, the strip is cooled on a run-out table to control the development of the final microstructure of the sheet prior to coiling into a steel roll.
While the three stage process of forming sheet in either twin roll casting or thin slab casting is part of the process, the response of the alloys herein to these stages is unique based on the mechanisms and structure types described herein and the resulting novel combinations of properties. Accordingly, in the present disclosure, sheet may be understood as metal formed into a relatively flat geometry of a selected thickness and width and slab may be understood as a length of metal that may be further processed into sheet material. Sheet may therefore be available as a relatively flat material or as a coiled stip.
Class 1 And Class 2 Steel
The alloys herein are such that they are capable of formation of what is described herein as Class 1 Steel or Class 2 Steel which are preferably crystalline (non-glassy) with identifiable crystalline grain size morphology. The ability of the alloys to form Class 1 or Class 2 Steels herein is described in detail herein. However, it is useful to first consider a description of the general features of Class 1 and Class 2 Steels, which is now provided below.
Class 1 Steel
The formation of Class 1 Steel herein is illustrated in FIG. 3A. As shown therein, a modal structure is initially formed which modal structure is the result of starting with a liquid melt of the alloy and solidifying by cooling, which provides nucleation and growth of particular phases having particular grain sizes. Reference herein to modal may therefore be understood as a structure having at least two grain size distributions. Grain size herein may be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Accordingly, Structure 1 of the Class 1 Steel may be preferably achieved by processing through either laboratory scale procedures as shown and/or through industrial scale methods involving chill surface processing methodology such as twin roll processing or thin slab casting
The modal structure of Class 1 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 500 nm to 20,000 nm containing austenite and/or ferrite; (2) boride grain size of 25 nm to 500 nm (i.e. non-metallic grains such as M2B where M is the metal and is covalently bonded to B). The boride grains may also preferably be “pinning” type phases which is reference to the feature that the matrix grains will effectively be stabilized by the pinning phases which resist coarsening at elevated temperature. Note that the metal boride grains have been identified as exhibiting the M2B stoichiometry but other stoichiometries are possible and may provide pinning including M3B, MB (M1B1), M23B6, and M7B3.
The modal structure of Class 1 Steel may be deformed by thermomechanical deformation and through heat treatment, resulting in some variation in properties, but the modal structure may be maintained.
When the Class 1 Steel noted above is exposed to a mechanical stress, the observed stress versus strain diagram is illustrated in FIG. 4. It is therefore observed that the modal structure undergoes what is identified as dynamic nanophase precipitation leading to a second type structure for the Class 1 Steel. Such dynamic nanophase precipitation is therefore triggered when the alloy experiences a yield under stress, and it has been found that the yield strength of Class 1 Steels which undergo dynamic nanophase precipitation may preferably occur at 300 MPa to 840 MPa. Accordingly, it may be appreciated that dynamic nanophase precipitation occurs due to the application of mechanical stress that exceeds such indicated yield strength. Dynamic nanophase precipitation itself may be understood as the formation of a further identifiable phase in the Class 1 Steel which is termed a precipitation phase with an associated grain size. That is, the result of such dynamic nanophase precipitation is to form an alloy which still indicates identifiable matrix grain size of 500 nm to 20,000 nm, boride pinning grain size of 25 nm to 500 nm, along with the formation of precipitation grains which contain hexagonal phases and grains of 1.0 nm to 200 nm. As noted above, the grain sizes therefore do not coarsen when the alloy is stressed, but does lead to the development of the precipitation grains as noted.
Reference to the hexagonal phases may be understood as a dihexagonal pyramidal class hexagonal phase with a P63mc space group (#186) and/or a ditrigonal dipyramidal class with a hexagonal P6bar2C space group (#190). In addition, the mechanical properties of such second type structure of the Class 1 Steel are such that the tensile strength is observed to fall in the range of 630 MPa to 1100 MPa, with an elongation of 10-40%. Furthermore, the second type structure of the Class 1 Steel is such that it exhibits a strain hardening coefficient between 0.1 to 0.4 that is nearly flat after undergoing the indicated yield. The strain hardening coefficient is reference to the value of n In the formula σ=Kεn, where σ represents the applied stress on the material, c is the strain and K is the strength coefficient. The value of the strain hardening exponent n lies between 0 and 1. A value of 0 means that the alloy is a perfectly plastic solid (i.e. the material undergoes non-reversible changes to applied force), while a value of 1 represents a 100% elastic solid (i.e. the material undergoes reversible changes to an applied force).
Table 1 below provides a comparison and performance summary for Class 1 Steel herein.







TABLE 1







Comparison of Structure and Performance for Class 1 Steel









Class 1 Steel









Property/
Structure Type #1
Structure Type #2 Modal 


Mechanism
Modal Structure
Nanophase Structure





Structure
Starting with a liquid melt,
Dynamic Nanophase Pre-


Formation
solidifying this liquid melt 
cipitation occurring 



and forming directly
through the application 




of mechanical stress


Transformations
Liquid solidification 
Stress induced 



followed by nucleation 
transformation involving 



and growth
phase formation




and precipitation


Enabling Phases
Austenite and/or ferrite 
Austenite, optionally 



with boride pinning
ferrite, boride pinning 




phases, and hexagonal




phase(s) precipitation


Matrix Grain
500 to 20,000 nm
500 to 20,000 nm


Size
Austenite and/or ferrite
Austenite optionally 




ferrite


Boride Grain Size
25 to 500 nm
25 to 500 nm



Non metallic 
Non-metallic 



(e.g. metal boride)
(e.g. metal boride)


Precipitation
—
1 nm to 200 nm


Grain Sizes

Hexagonal phase(s)


Tensile Response
Intermediate structure; 
Actual with properties 



transforms into Structure #2 
achieved based 



when undergoing yield
on structure type #2


Yield Strength
300 to 600 MPa
300 to 840 MPa


Tensile Strength
—
630 to 1100 MPa


Total Elongation
—
10 to 40%


Strain Hardening
—
Exhibits a strain 


Response

hardening coefficient 




between 0.1 to 0.4 




and a strain hardening




coefficient as a function




of strain which is nearly




flat or experiencing a 




slow increase until failure









Class 2 Steel

As shown in FIG. 3B, Class 2 steel may also be formed herein from the identified alloys, which unlike Class 1 Steel, involves two new structure types after starting with Structure type #1 of Class 1 Steel, but followed by two new mechanisms identified herein as static nanophase refinement and dynamic nanophase strengthening. The new structure types for Class 2 Steel are described herein as nanomodal structure and high strength nanomodal structure. Accordingly, Class 2 Steel herein may be characterized as follow: Structure #1—Modal Structure (Step #1), Mechanism #1—Static Nanophase Refinement (Step #2), Structure #2—NanoModal Structure (Step #3), Mechanism #2—Dynamic Nanophase Strengthening (Step #4), and Structure #3—High Strength NanoModal Structure (Step #5). Structure #1 involving the formation of the modal structure in the Class 2 Steel is the same as for Class 1 Steel above and may again be achieved in the alloys with the referenced chemistries in this application by processing through either laboratory scale procedures as disclosed herein and/or through industrial scale methods involving chill surface processing methodology such as twin roll processing or thin slab casting. Reference to Structure 1—Modal Structure of Class 2 Steel herein may therefore again be understood as having grain sizes in the range of 500 nm to 20,000 nm and an identifiable boride grain size of 25 nm to 500 nm (which is metal boride grain phase such as exhibiting the M2B stoichiometry or also other stoichiometries such as M3B, MB (M1B1), M23B6, and M7B3, and which is unaffected by mechanism 1 or 2 noted above). Reference to grain size is again to be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Furthermore, Structure 1 of Class 2 steel herein includes austenite and/or ferrite along with such boride phases. In addition the boride phase, as in Class 1 Steel is preferably a pinning phase.
In FIG. 5, a stress strain curve is shown that represents the alloys herein which undergo a deformation behavior of a representative Class 2 steel. The modal structure is again preferably first created (Structure #1) and then after the creation, the modal structure may now be refined (i.e. the grain size distribution is altered) through Mechanism #1, which is a Static Nanophase Refinement mechanism, leading to Structure #2. Static Nanophase Refinement is reference to the feature that the matrix grain sizes of Structure 1 which initially fall in the range of 500 nm to 20,000 nm are reduced in size to provide Structure 2 which has matrix grain sizes that typically fall in the range of 100 nm to 2000 nm. Note that the boride pinning phase does not change significantly in size and thus resists coarsening during the heat treatments. Due to the presence of these boride pinning sites, the motion of a grain boundaries leading to coarsening would be expected to be retarded by a process called Zener pinning or Zener drag. The boride phases which are non-metallic would exhibit a high interfacial energy which is lowered by existing at grain or phase boundaries. Thus, while grain growth of the matrix may be energetically favorable due to the reduction of total interfacial area, the presence of the boride pinning phase will counteract this driving force of coarsening due to the high interfacial energies of these phases. Structure 2 also displays completely different behavior when tested in tension and has the potential to achieve much higher strengths than a Class 1 Steel.
Characteristic of the Static Nanophase Refinement mechanism in Class 2 steel, the micron scale austenite phase (gamma-Fe) which was noted as falling in the range of 500 nm to 20,000 nm is partially or completely transformed into new phases (e.g. ferrite or alpha-Fe). The volume fraction of ferrite initially present in the modal structure of Class 2 steel is 10 to 70%. The volume fraction of ferrite (alpha-iron) in Structure 2 as a result of Static Nanophase Refinement is typically from 20 to 80%. The static transformation preferably occurs during elevated temperature heat treatment and thus involves a unique refinement mechanism since grain coarsening and not grain refinement is the conventional material response at elevated temperature. Accordingly, grain coarsening does not occur with the alloys of Class 2 Steel herein during the Static Nanophase Refinement mechanism. Structure 2 is uniquely able to transform to Structure #3 during Dynamic Nanophase Strengthening and as a result Structure#3 is formed and indicates tensile strength values in the range from 720 to 1580 MPa tensile strength and 5 to 35% total elongation.
Expanding upon the above, in the case of the alloys herein that provide Class 2 Steel, when such alloys exceed their yield point, plastic deformation at constant stress occurs followed by a dynamic phase transformation leading toward the creation of Structure #3. More specifically, after enough strain is induced, an inflection point occurs where the slope of the stress versus strain curve changes and increases (FIG. 5) and the strength increases with strain indicating an activation of Mechanism #2 (Dynamic Nanophase Strengthening). An increase in strain hardening coefficient is also found at the beginning of deformation. The value of the strain hardening exponent n lies between 0.2 to 1.0 for Structure 3 in the Class 2 Steel.
With further straining during Dynamic Nanophase Strengthening, the strength continues to increase but with a gradual decrease in strain hardening coefficient value up to nearly failure. Some strain softening occurs near the breaking point which may be due to reductions in localized cross sectional area at necking. Note that the strengthening transformation that occurs at the material straining under the stress generally defines Mechanism #2 as a dynamic process, leading to Structure #3. By dynamic, it is meant that the process may occur through the application of a stress which exceeds the yield strength of the material. The tensile properties that can be achieved for alloys that achieve Structure 3 include tensile strength values in the range from 720 to 1580 MPa tensile strength and 5 to 35% total elongation. The level of tensile properties achieved is also dependant on the amount of transformation occurring as the strain is increased corresponding to the characteristic stress strain curve for a Class 2 steel.
Thus, depending on the level of transformation, a tunable yield strength may also now be developed in Class 2 Steel herein depending on the level of deformation and in Structure 3 the yield strength can ultimately vary from 300 MPa to 1300 MPa. That is, conventional steels outside the scope of the alloys here exhibit only relatively low levels of strain hardening, thus their yield strengths can be varied only over small ranges (e.g., 100 to 200 MPa) depending on the prior deformation history. In Class 2 steels herein, the yield strength can be varied over a wide range (e.g. 300 to 600 MPa) as applied to Structure 2, allowing tunable variations to enable both the designer and end users in a variety of applications to achieve Structure 3, and utilize Structure 3 in various applications such as crash management in automobile body structures.
With regards to this dynamic mechanism shown in FIG. 3B, a new precipitation phase is observed that indicates identifiable grain sizes of 1 nm to 200 nm. In addition, there is the further identification in said precipitation phase a dihexagonal pyramidal class hexagonal phase with a P63mc space group (#186) and/or a ditrigonal dipyramidal class with a hexagonal P6bar2C space group (#190). Accordingly, the dynamic transformation can occur partially or completely and results in the formation of a microstructure with novel nanoscale/near nanoscale phases providing relatively high strength in the material. That is, Structure #3 may be understood as a microstructure having a matrix grain size generally from 100 nm to 2000 nm which are pinned by boride phases which are in the range of 25 to 500 nm and with precipitate phases which are in the range of 1 nm to 200 nm.
Note that dynamic recrystallization is a known process but differs from Mechanism #2 since it involves the formation of large grains from small grains so that it is not a refinement mechanism but a coarsening mechanism. Additionally, as new undeformed grains are replaced by deformed grains no phase changes occur in contrast to the mechanisms presented here and this also results in a corresponding reduction in strength in contrast to the strengthening mechanism here. Note also that metastable austenite in steels is known to transform to martensite under mechanical stress but, preferably, no evidence for martensite or body centered tetragonal iron phases are found in the new steel alloys described in this application. Table 2 below provides a comparison of the structure and performance features of Class 2 Steel herein.







TABLE 2







Comparison Of Structure and Performance of Class 2 Steel









Class 2 Steel













Structure Type #3


Property/
Structure Type #1
Structure Type #2
High Strength


Mechanism
Modal Structure
NanoModal Structure
NanoModal Structure





Structure
Starting with a liquid melt,
Static Nanophase Refinement
Dynamic Nanophase


Formation
solidifying this liquid melt
mechanism occurring during
Strengthening mechanism



and forming directly
heat treatment
occurring through





application of mechanical





stress


Transformations
Liquid solidification
Solid state phase
Stress induced



followed by nucleation and
transformation of
transformation involving



growth
supersaturated gamma iron
phase formation and





precipitation


Enabling Phases
Austenite and/or ferrite with
Ferrite, austenite, boride
Ferrite, optionally austenite,



boride pinning phases
pinning phases
boride pinning phases, and





hexagonal phase(s)





precipitation


Matrix Grain
500 to 20,000 nm
Grain Refinement
Grain size remains refined


Size
Austenite and/or ferrite
(100 nm to 2000 nm)
at 100 nm to 2000 nm/




Austenite phase to ferrite
Hexagonal phase formation




phase



Boride Grain Size
25 to 500 nm
25 to 500 nm
25 to 500 nm



borides (e.g. metal boride)
borides (e.g. metal boride)
borides (e.g. metal boride)


Precipitation
—
—
1 nm to 200 nm


Grain Sizes


Hexagonal phase(s)


Tensile Response
Actual with properties
Intermediate structure;
Actual with properties



achieved based on structure
transforms into Structure #3
achieved based on



type #1
when undergoing yield
formation of structure type





#3 and fraction of





transformation.


Yield Strength
300 to 600 MPa
300 to 600 MPa
300 to 1300 MPa


Tensile Strength
—
—
720 to 1580 MPa


Total Elongation
—
—
5 to 35%


Strain Hardening
—
After yield point, exhibit a
Strain hardening coefficient


Response

strain softening at initial
may vary from 0.2 to 1.0




straining as a result of phase
depending on amount of




transformation, followed by a
deformation and




significant strain hardening
transformation




affect leading to a distinct





maxima









Mechanisms During Production
The formation of Modal Structure (MS) in either Class 1 or Class 2 Steel herein can be made to occur at various stages of the production process. For example, the MS of the sheet may form during Stage 1, 2, or 3 of either the above referenced twin roll or thin slab casting sheet production processes. Accordingly, the formation of MS may depend specifically on the solidification sequence and thermal cycles (i.e. temperatures and times) that the sheet is exposed to during the production process. The MS may be preferably formed by heating the alloys herein at temperatures in the range of above their melting point and in a range of 1100° C. to 2000° C. and cooling below the melting temperature of the alloy, which corresponds to preferably cooling in the range of 11×103 to 4×10−2K/s.
With respect to Class 2 Steel herein, Mechanism #1 which is the Static Nanophase Refinement (SNR) occurs after MS is formed and during further elevated temperature exposure. Accordingly, Static Nanophase Refinement may also occur during Stage 1, Stage 2 or Stage 3 (after MS formation) of either of the above referenced twin roll or thin slab casting sheet production process. It has been observed that Static Nanophase Refinement may preferably occur when the alloys are subject to heating at temperature in the range of 700° C. to 1200° C. The percentage level of SNR that occurs in the material may depend on the specific chemistry and involved thermal cycle that determines the volume fraction of NanoModal Structure (NMS) specified as Structure #2. However, preferably, the percentage level by volume of MS that is converted to NMS is in the range of 20 to 90%.
Mechanism #2 which is Dynamic Nanophase Strengthening (DNS) may also occur during Stage 1, Stage 2 or Stage 3 (after MS formation) of either of the above referenced twin roll or thin slab casting sheet production process. Dynamic Nanophase Strengthening may therefore occur in Class 2 Steel that has undergone Static Nanophase Refinement. Dynamic Nanophase Strengthening may therefore also occur during the production process of sheet but may also be done during any stage of post processing involving application of stresses exceeding the yield strength. Tables 6 and 8 relate to tensile measurements where Dynamic NanoPhase Strengthening is occurring since the heat treatment(s) caused the creation of the NanoModal Structure. The amount of DNS that occurs may depend on the volume fraction of static nanophase refinement in the material prior deformation and on stress level induced in the sheet. The strengthening may also occur during subsequent post processing into final parts involving hot or cold forming of the sheet. Thus Structure #3 herein (see Table 2 above) may occur at various processing stages in the sheet production or upon post processing and additionally may occur to different levels of strengthening depending on the alloy chemistry, deformation parameters and thermal cycle(s). Preferably, DNS may occur under the following range of conditions, after achieving structure type #2 and then exceeding the yield strength of the structure which is in the range of 300 to 1300 MPa.
FIG. 3C illustrates in general that starting with a particular chemical composition for the alloys herein, and heating to a liquid, and solidifying on a chill surface, and forming modal structure, one may then convert to either Class 1 Steel or Class 2 Steel as noted herein.
EXAMPLES
Preferred Alloy Chemistries and Sample Preparation
The chemical composition of the alloys studied is shown in Table 2 which provides the preferred atomic ratios utilized. These chemistries have been used for material processing through sheet casting in a Pressure Vacuum Caster (PVC). Using high purity elements [>99 wt %], 35 g alloy feedstocks of the targeted alloys were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a PVC chamber, melted using RF induction and then ejected onto a copper die designed for casting 3 by 4 inches sheets with thickness of 1.8 mm.







TABLE 2







Chemical Composition of the Alloys

















Alloy
Fe
Cr
Ni
B
Si
V
Zr
C
W
Mn




















Alloy 1
59.35
17.43
14.05
4.77
4.40
—
—
—
—



Alloy 2
57.75
17.43
14.05
4.77
6.00
—
—
—
—



Alloy 3
58.35
17.43
14.05
4.77
4.40
1.00
—
—
—



Alloy 4
54.52
17.43
14.05
7.00
7.00
—
—
—
—



Alloy 5
56.52
17.43
14.05
7.00
5.00
—
—
—
—



Alloy 6
55.52
17.43
14.05
7.00
5.00
1.00
—
—
—



Alloy 7
53.52
17.43
14.05
7.00
5.00
3.00
—
—
—



Alloy 8
53.52
17.43
14.05
7.00
7.00
1.00
—
—
—



Alloy 9
55.52
17.43
14.05
7.00
5.00
—
1.00
—
—



Alloy 10
57.35
17.43
14.05
4.77
4.40
—
—
2.00
—



Alloy 11
66.35
17.43
7.05
4.77
4.40
—
—
—
—



Alloy 12
58.35
17.43
14.05
4.77
4.40
—
—
—
1.00



Alloy 13
57.22
17.43
14.05
5.00
6.30
—
—
—
—



Alloy 14
64.22
17.43
7.05
5.00
6.30
—
—
—
—



Alloy 15
63.22
17.43
7.05
5.00
6.30
—
—
—
1.00



Alloy 16
68.70
15.00
5.00
5.00
6.30
—
—
—
—



Alloy 17
64.75
17.43
7.05
4.77
6.00
—
—
—
—



Alloy 18
65.45
17.43
9.05
4.47
5.60
—
—
—
—



Alloy 19
63.62
17.43
12.05
5.30
6.60
—
—
—
—



Alloy 20
62.22
17.43
9.05
5.00
6.30
—
—
—
—



Alloy 21
60.22
17.43
11.05
5.00
6.30
—
—
—
—



Alloy 22
62.22
19.43
7.05
5.00
6.30
—
—
—
—



Alloy 23
66.22
15.43
7.05
5.00
6.30
—
—
—
—



Alloy 24
62.75
17.43
9.05
4.77
6.00
—
—
—
—



Alloy 25
62.20
17.62
4.14
5.30
6.60




4.14


Alloy 26
60.35
20.70
3.53
5.30
6.60




3.52


Alloy 27
61.10
19.21
3.90
5.30
6.60




3.89


Alloy 28
61.32
20.13
3.33
5.30
6.60




3.32


Alloy 29
63.83
17.97
3.15
5.30
6.60




3.15


Alloy 30
63.08
15.95
4.54
5.30
6.60




4.53


Alloy 31
64.93
16.92
3.13
5.30
6.60




3.12


Alloy 32
64.45
15.86
3.90
5.30
6.60




3.89


Alloy 33
62.11
20.31
2.84
5.30
6.60




2.84


Alloy 34
62.20
17.62
6.21
5.30
6.60




2.07


Alloy 35
60.35
20.70
5.29
5.30
6.60




1.76


Alloy 36
61.10
19.21
5.85
5.30
6.60




1.94


Alloy 37
61.32
20.13
4.99
5.30
6.60




1.66


Alloy 38
63.83
17.97
4.73
5.30
6.60




1.57


Alloy 39
63.08
15.95
6.80
5.30
6.60




2.27


Alloy 40
64.93
16.92
4.69
5.30
6.60




1.56


Alloy 41
64.45
15.86
5.85
5.30
6.60




1.94


Alloy 42
62.11
20.31
4.26
5.30
6.60




1.42


Alloy 43
72.10
12.20
4.50
7.20
4.00







Alloy 44
62.38
17.40
7.92
7.40
4.20


0.20

0.50


Alloy 45
65.99
13.58
6.58
7.60
4.40


0.35

1.50


Alloy 46
58.76
17.22
9.77
7.80
4.60


0.55

1.30


Alloy 47
58.95
11.35
13.40
8.00
4.80


2.25

1.25


Alloy 48
62.28
10.00
12.56
4.80
8.00


0.36

2.00


Alloy 49
53.82
20.22
11.60
4.60
7.80


1.21

0.75


Alloy 50
61.21
21.00
4.90
4.40
7.60


0.89




Alloy 51
62.00
17.50
6.25
4.20
7.40


2.55

0.10


Alloy 52
59.71
14.30
13.74
4.00
7.20


0.65

0.40


Alloy 53
57.85
13.90
12.25
7.00
7.00


0.25

1.75


Alloy 54
56.90
15.25
14.50
6.00
6.00




1.35


Alloy 55
65.82
12.22
7.22
5.00
6.00


2.60

1.14


Alloy 56
58.72
18.26
8.99
4.26
7.22


1.00

1.55


Alloy 57
61.30
17.30
6.50
7.15
4.55


3.00

0.20


Alloy 58
65.80
14.89
8.66
4.35
4.05


2.25




Alloy 59
63.99
12.89
10.25
8.00
4.22




0.65


Alloy 60
71.24
10.55
5.22 
7.55
4.55




0.89


Alloy 61
61.88
11.22
12.55
7.45
5.22


0.56

1.12









Accordingly, in the broad context of the present disclosure, the alloy chemistries that may preferably be suitable for formation of the Class 1 or Class 2 Steel herein include the following elements whose atomic ratios add up to 100. That is, the alloys may include Fe, Cr, Ni, B and Si. The alloys may optionally include V, Zr, C, W or Mn. Preferably, with respect to atomic ratios, the alloys may contain Fe at 53.5 to 72.1, Cr at 10.0 to 21.0, Ni at 2.8 to 14.50, B at 4.00 to 8.00 and Si at 4.00 to 8.00, and optionally V at 1.0 to 3.0, Zr at 1.00, C at 0.2 to 3.00, W at 1.00, or Mn at 0.20 to 4.6. Accordingly, the levels of the particular elements may be adjusted to total 100 as noted above.
The atomic ratio of Fe present may therefore be 53.5, 53.6, 53.7, 54.8, 53.9, 53.0 53.1, 53.2, 53.3, 53.4, 53.5, 53.6, 53.7, 53.8, 53.9, 54.0, 54.1, 54.2, 54.3, 54.4, 54.5, 54.6, 54.7, 54.8, 54.9, 55.0, 55.1, 55.2, 55.3, 55.4, 55.5, 55.6, 55.7, 55.8, 55.9, 56.0, 56.1, 56.2, 56.3, 56.4, 56.5, 56.6, 56.7, 56.8, 56.9 57.0, 57.1, 57.2, 57.3, 57.4, 57.5, 57.6, 57.7, 57.8, 57.9, 58.0, 58.1, 58.2, 58.3, 58.4, 58.5, 58.6, 58.7, 58.8, 58.9, 59.0, 59.1, 59.2, 59.3, 59.4, 59.5, 59.6, 59.7, 59.8, 60.0, 60.1, 60.2, 60.3, 60.4, 60.5, 60.6, 60.7, 60.8, 60.9 61.0, 61.1, 61.2, 61.3, 61.4, 61.5, 61.6, 61.7, 61.8, 61.9, 62.0, 62.1, 62.2, 62.3, 62.4, 62.5, 62.6, 62.7, 62.8, 62.9, 63.0, 63.1, 63.2, 63.3, 63.4, 63.5, 63.6, 63.7, 63.8, 63.9, 64.0, 64.1, 64.2, 64.3, 64.4, 64.5, 64.6, 64.7, 64.8, 64.9, 65.0, 65.1, 65.2, 65.3, 65.4, 65.5, 65.6, 65.7, 65.8, 65.9, 66.0, 66.1, 66.2, 66.3, 66.4, 66.5, 66.6, 66.7, 66.8, 66.9, 67.0, 67.1, 67.2, 67.3, 67.4, 67.5, 67.6, 67.7, 67.8, 67.9, 68.0, 68.1, 68.2, 68.3, 68.4, 68.5, 68.6, 68.7, 68.8, 68.9, 69.0, 70.0, 70.1, 70.2, 70.3, 70.4, 70.5, 70.6, 70.7, 70.8, 70.9, 71.0, 71.1, 71.2, 71.3, 71.4, 71.5, 71.6, 71.7, 71.8, 71.9, 72.0, 72.1. The atomic ratio of Cr may therefore be 10.0, 10.1, 10.2, 10.3, 10.4, 10.5, 10.6, 10.7, 10.8, 10.9, 11.0, 11.1, 11.2, 11.3, 11.4, 11.5, 11.6, 11.7, 11.8, 11.9, 12.0, 12.1, 12.2, 12.3, 12.4, 12.5, 12.6, 12.7, 12.8, 12.9 13.0, 13.1, 13.2, 13.3, 13.4, 13.5, 13.6, 13.7, 13.8, 13.9, 14.0, 14.1, 14.2, 14.3, 14.4, 14.5, 14.6, 14.7, 14.8, 14.9, 15.0, 15.1, 15.2, 15.3, 15.4, 15.5, 15.6, 15.7, 15.8, 15.9, 16.0, 16.1, 16.2, 16.3, 16.4, 16.5, 16.6, 16.7, 16.8, 16.9, 17.0, 17.1, 17.2, 17.3, 17.4, 17.5, 17.6, 17.7, 17.8, 17.9, 18.0, 18.1, 18.2, 18.3, 18.4, 18.5, 18.6, 18.7, 18.8, 18.9, 19.0, 19.1, 19.2, 19.3, 19.4, 19.5, 19.6, 19.7, 19.8, 19.9, 20.0, 20.1, 20.2, 20.3, 20.4, 20.5, 20.6, 20.7, 20.8, 20.9, 21.0. The atomic ratio of Ni may therefore be 2.8, 2.9 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6, 5.7, 5.8, 5.9, 6.0, 6.1, 6.2, 6.3, 6.4, 6.5, 6.6, 6.7, 6.8, 6.9, 7.0, 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9, 8.0, 8.1, 8.2, 8.3, 8.4, 8.5, 8.6, 8.7, 8.8, 8.9, 9.0, 9.1, 9.2, 9.3, 9.4, 9.5, 9.6, 9.7, 9.8, 9.9 10.0, 10.1, 10.2, 10.3, 10.4, 10.5, 10.6, 10.7, 10.8, 10.9, 11.0, 11.1, 11.2, 11.3, 11.4, 11.5, 11.6, 11.7, 11.8, 11.9, 12.0, 12.1, 12.2, 12.3, 12.4, 12.5, 12.6, 12.7, 12.8, 12.9, 13.0, 13.1, 13.2, 13.3, 13.4, 13.5, 13.6, 13.7, 13.8, 13.9, 14.0, 14.1, 14.2, 14.3, 14.4, 14.50. The atomic ratio of B may therefore be 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6, 5.7, 5.8, 5.9, 6.0, 6.1, 6.2, 6.3, 6.4, 6.5, 6.6, 6.7, 6.8, 6.9, 7.0, 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9, 8.0. The atomic ratio of Si may therefore be 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6, 5.7, 5.8, 5.9, 6.0, 6.1, 6.2, 6.3, 6.4, 6.5, 6.6, 6.7, 6.8, 6.9, 7.0, 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9, 8.0. The atomic ratio of Si may therefore be 4.0, 5.0, 6.0, 7.0, 8.0. The atomic ratio of the optional elements such as V may therefore be 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0. The atomic ratio of C may therefore be 0.2, 0.3, 0.4, 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0. The atomic ratio of W may therefore be 1.0. The atomic ratio of Mn may therefore be 0.20, 0.3, 0.4, 0.5, 0.6, 0.7, 0.8, 0.9 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6.
The alloys may herein may also be more broadly described as an Fe based alloy (greater than or equal to 50.00 atomic percent) and including B and Si at levels of 4.00 atomic percent to 8.00 atomic percent and capable of forming the indicated structures (Class 1 and/or Class 2 Steel) and/or undergoing the indicated transformations upon exposure to mechanical stress and/or mechanical stress in the presence of heat treatment. Such alloys may be further defined by the mechanical properties that are achieved for the identified structures with respect to tensile strength and tensile elongation characteristics.
Alloy Properties
Thermal analysis was done on the as-solidified cast sheet samples on a NETZSCH DSC 404F3 PEGASUS V5 system. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) was performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultrahigh purity argon. In Table 3, elevated temperature DTA results are shown indicating the melting behavior for the alloys. As can be seen from the tabulated results in Table 3, the melting occurs in 1 to 3 stages with initial melting observed from  ̃1184° C. depending on alloy chemistry. Final melting temperature is up to  ̃1340° C. Variations in melting behavior may also reflect a complex phase formation at chill surface processing of the alloys depending on their chemistry.







TABLE 3







Differential Thermal Analysis Data for Melting Behavior












Onset
Peak #1
Peak #2
Peak #3


Alloy
(° C.) 
(° C.)
(° C.)
(° C.)





Alloy 1
1234
1258
1331
—


Alloy 2
1233
1252
1318
—


Alloy 3
1230
1254
1325
—


Alloy 4
1187
1233
—
—


Alloy 5
1204
1246
1268
—


Alloy 6
1203
1241
—
—


Alloy 7
1207
1237
—
—


Alloy 8
1184
1232
—
—


Alloy 9
1190
1203
1235
—


Alloy 10
1188
1195
1246
1314


Alloy 11
1243
1256
1345
—


Alloy 12
1221
1248
1330
—


Alloy 13
1221
1248
1305
—


Alloy 14
1231
1251
1330
—


Alloy 15
1225
1241
1321
—


Alloy 16
1225
1241
1338
—


Alloy 17
1227
1245
1335
—


Alloy 18
1225
1244
1340
—


Alloy 19
1222
1239
1309
—


Alloy 20
1221
1245
1309
—


Alloy 21
1209
1242
1299
—


Alloy 22
1223
1250
1315
—


Alloy 23
1209
1234
1316
—


Alloy 24
1222
1241
1316
—









The density of the alloys was measured on arc-melt ingots using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water. The density of each alloy is tabulated in Table 4 and was found to vary from 7.53 g/cm3 to 7.77 g/cm3. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3.







TABLE 4







Summary of Density Results (g/cm3)











Density 



Alloy
(avg)







Alloy 1 
7.73



Alloy 2 
7.68



Alloy 3 
7.73



Alloy 4 
7.60



Alloy 5 
7.65



Alloy 6 
7.64



Alloy 7 
7.60



Alloy 8 
7.57



Alloy 9 
7.66



Alloy 10
7.70



Alloy 11
7.63



Alloy 12
7.91



Alloy 13
7.67



Alloy 14
7.61



Alloy 15
7.77



Alloy 16
7.49



Alloy 17
7.62



Alloy 18
7.64



Alloy 19
7.58



Alloy 20
7.64



Alloy 21
7.65



Alloy 22
7.60



Alloy 23
7.53



Alloy 24
7.65










The tensile specimens were cut from the sheets using wire electrical discharge machining (EDM). The tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving; the load cell is attached to the top fixture. In Table 5, a summary of the tensile test results including total tensile strain, yield stress, ultimate tensile strength, Elastic Modulus and strain hardening exponent value are shown for as-cast sheets. The mechanical characteristic values depend on alloy chemistry and processing condition as will be discussed herein. As can be seen the ultimate tensile strength values vary from 590 to 1290 MPa. The tensile elongation varies from 0.79 to 11.27%. Elastic Modulus is measured in a range from 127 to 283 GPa. Strain hardening coefficient was calculated in a range from 0.13 to 0.44







TABLE 5







Summary on Tensile Test Results for As-Cast Sheets















Ultimate
Tensile






Yield 
Tensile
Elon- 
Elastic
Strain
Type 



Stress
Strength
gation
Modulus
Hardening 
of



(MPa)
(MPa)
(%)
(GPa)
Exponent
Behavior





Alloy 1
430
830
4.66
177
0.28
Class 1



490
720
2.63
175
0.23
Class 1



440
770
5.87
163
0.23
Class 1


Alloy 2
500
810
4.06
161
0.25
Class 1



400
840
3.71
165
0.27
Class 1



500
770
5.29
172
0.23
Class 1



400
840
6.10
169
0.27
Class 1


Alloy 3
500
950
9.77
156
0.24
Class 1



500
900
6.49
171
0.25
Class 1



500
920
10.53 
181
0.25
Class 1



400
890
11.27 
177
0.24
Class 1


Alloy 4
590
960
2.53
173
0.29
Class 1



600
970
2.77
185
0.29
Class 1



600
710
0.79
197
0.32
Class 1


Alloy 5
480
840
1.74
162
0.31
Class 1



620
1010 
3.34
190
0.26
Class 1



600
910
2.45
205
0.25
Class 1



540
760
1.43
160
0.32
Class 1


Alloy 6
570
810
1.57
191
N/A
Class 1



580
930
2.45
189
0.28
Class 1



620
1030 
2.99
201
0.26
Class 1


Alloy 7
560
860
1.86
178
0.28
Class 1



530
730
1.01
283
N/A
Class 1



560
940
2.85
187
0.28
Class 1


Alloy 8
600
930
2.20
182
0.29
Class 1



620
760
0.97
190
0.32
Class 1


Alloy 9
430
640
1.30
144
N/A
Class 1


Alloy 10
560
1030 
3.56
184
0.31
Class 1


Alloy 11
500
890
5.83
172
0.23
Class 1



500
820
5.83
180
0.19
Class 1


Alloy 12
430
870
8.35
172
0.27
Class 1



390
590
1.97
172
0.28
Class 1


Alloy 13
470
800
3.73
170
0.26
Class 1



410
720
2.32
185
0.31
Class 1


Alloy 14
670
840
1.19
178
N/A
Class 1


Alloy 15
690
930
1.87
164
0.24
Class 1


Alloy 16
770
1010 
1.06
186
0.44
Class 2



900
1290 
1.56
185
0.44
Class 2


Alloy 17
590
780
1.30
203
N/A
Class 1



710
820
1.02
196
N/A
Class 1



670
820
1.20
181
N/A
Class 1



650
860
2.02
243
0.15
Class 1


Alloy 18
540
830
5.24
127
0.15
Class 1



560
1010 
7.93
164
0.23
Class 1



550
940
7.36
168
0.19
Class 1



570
840
5.14
178
0.13
Class 1



570
850
5.84
177
0.15
Class 1



660
1020 
7.07
174
0.18
Class 1


Alloy 19
670
910
1.90
181
0.23
Class 1



630
840
1.41
161
N/A
Class 1



620
730
1.02
155
N/A
Class 1



610
960
2.34
212
0.27
Class 1



760
990
2.09
202
0.18
Class 1


Alloy 20
540
1040 
6.23
193
0.26
Class 1



560
1040 
6.85
195
0.23
Class 1



520
850
2.59
174
0.29
Class 1



460
890
3.25
173
0.29
Class 1


Alloy 21
450
880
6.69
148
0.27
Class 1



450
850
2.96
200
0.30
Class 1



450
770
2.72
175
0.30
Class 1



410
640
1.98
163
0.30
Class 1


Alloy 22
600
800
1.19
191
N/A
Class 1



840
1060 
2.15
140
0.24
Class 1



750
1100 
2.30
181
0.25
Class 1



730
1000 
1.99
178
0.25
Class 1


Alloy 23
420
810
2.82
148
0.36
Class 1



410
700
2.80
146
0.30
Class 1


Alloy 24
490
850
3.05
180
0.27
Class 1



510
970
6.87
184
0.23
Class 1









Alloy Properties after Thermal Mechanical Treatment
Each sheet from each alloy was subjected to Hot Isostatic Pressing (HIP) using an American Isostatic Press Model 645 machine with a molybdenum furnace and with a furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time which was held at 1 hour for these studies. HIP cycle parameters are listed in Table 6. The preferred aspect of the HIP cycle was to remove macrodefects such as pores (0.5 to 100 μm) and small inclusions (0.5 to 100 μm) by mimicking hot rolling at Stage 2 of Twin Roll Casting process or at Stage 1 or Stage 2 of Thin Slab Casting process. An example sheet before and after HIP cycle is shown in FIG. 6. As it can be seen, the HIP cycle which is a thermomechanical deformation process allows the elimination of some fraction of internal and external macrodefects and smoothes the surface of the sheet.







TABLE 6







HIP Cycle Parameters












HIP
HIP Cycle
HIP Cycle
HIP Cycle



Cycle 
Temperature
Pressure
Time



ID
[° C.]
[psi]
[hr]







Ha
 700
30,000
1



Hb
 850
30,000
1



Hd
 900
30,000
1



Hc
1000
30,000
1



He
1100
30,000
1



Hf
1150
30,000
1










The tensile specimens were cut from the sheets after HIPing using wire electrical discharge machining (EDM). The tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving with the load cell attached to the top fixture. In Table 7, a summary of the tensile test results including total tensile strain, yield stress, ultimate tensile strength, Elastic Modulus and strain hardening exponent value are shown for the cast sheets after HIP cycle. Mechanical characteristic values strongly depend on alloy chemistry and HIP cycle parameters. As can be seen the ultimate tensile strength values vary from 630 to 1440 MPa. The tensile elongation value varies from 1.11 to 24.41%. Elastic Modulus was measured in a range from 121 to 230 GPa. Strain hardening coefficient was calculated from the yield strength to the tensile strength resulting in ranges from 0.13 to 0.99 depending on alloy chemistry, structural formation, and different heat treatments.







TABLE 7







Summary on Tensile Test Results for HIPed Sheets

















Ultimate
Tensile






HIP
Yield
Tensile
Elon-
Elastic
Strain




Cycle
Stress
Strength
gation
Modulus
Hardening
Type of


Alloy
ID
(MPa)
(MPa)
(%)
(GPa)
Exponent
Behavior

















Alloy 1
Ha
460
870
4.12
163
0.27
Class 1




460
990
10.82
186
0.25
Class 1



Hb
400
750
5.10
147
0.28
Class 1




410
770
5.03
173
0.27
Class 1




400
800
6.79
132
N/A
Class 1




380
690
4.25
147
0.27
Class 1



Hc
340
790
14.64
170
0.27
Class 1




370
850
18.46
160
0.29
Class 1


Alloy 2
Ha
410
800
5.80
162
N/A
Class 1




410
860
7.99
142
0.27
Class 1



Hb
400
850
5.76
173
0.27
Class 1




500
910
9.17
165
0.25
Class 1




500
910
8.28
192
0.24
Class 1



Hc
400
910
21.16
168
0.25
Class 1




400
900
19.65
190
0.25
Class 1


Alloy 3
Ha
450
920
6.54
166
0.27
Class 1




450
950
8.37
181
0.25
Class 1



Hb
420
890
17.77
164
0.25
Class 1




430
920
12.24
172
0.26
Class 1



Hc
380
790
8.49
160
0.26
Class 1




360
790
13.40
194
0.26
Class 1


Alloy 4
Ha
610
1000
3.00
174
0.29
Class 1




600
950
2.04
187
0.31
Class 1



Hb
510
830
1.80
183
0.34
Class 1




560
870
2.11
177
0.31
Class 1



Hc
470
940
7.13
167
0.27
Class 1




460
970
9.35
168
0.27
Class 1


Alloy 5
Ha
580
970
2.75
180
0.29
Class 1




580
950
2.85
171
0.28
Class 1



Hb
510
970
4.32
208
0.27
Class 1




560
910
3.26
155
0.29
Class 1



Hc
470
970
10.06
177
0.25
Class 1




470
950
8.36
212
0.25
Class 1


Alloy 6
Ha
600
990
2.99
177
0.28
Class 1




570
900
2.17
183
0.30
Class 1



Hb
580
1000
3.51
184
0.28
Class 1




540
880
2.29
169
0.30
Class 1



Hc
490
930
5.81
184
0.27
Class 1




490
970
8.89
191
0.25
Class 1




470
910
5.01
179
0.28
Class 1


Alloy 7
Ha
590
810
1.16
196
N/A
Class 1




590
970
2.43
193
0.29
Class 1



Hb
580
970
2.95
176
0.29
Class 1




600
790
1.11
180
N/A
Class 1




560
1010
3.89
176
0.29
Class 1



Hc
470
820
2.7S
175
0.31
Class 1




480
890
4.42
175
0.27
Class 1


Alloy 8
Ha
590
1030
2.86
186
0.31
Class 1



Hb
570
1020
3.17
177
0.30
Class 1



Hc
490
860
3.13
192
0.30
Class 1




500
780
2.20
190
0.28
Class 1




530
860
2.86
173
0.30
Class 1


Alloy 10
Hb
530
1030
4.47
180
0.31
Class 1




530
1010
4.36
167
0.31
Class 1


Alloy 11
Hb
410
800
4.02
179
0.49
Class 2




410
950
4.71
194
0.76
Class 2



Hc
540
1060
2.13
174
0.51
Class 2




510
1330
7.97
133
0.43
Class 2




520
1320
7.39
169
0.35
Class 2


Alloy 12
Ha
430
770
2.87
131
0.29
Class 1




450
890
7.05
121
0.28
Class 1



Hb
440
890
5.51
159
0.28
Class 1




450
870
5.02
170
0.28
Class 1



Hc
400
870
12.73
177
0.24
Class 1




440
880
12.88
145
0.24
Class 1


Alloy 13
Hb
460
850
5.13
149
0.27
Class 1




380
820
5.57
154
0.30
Class 1



Hc
420
860
9.95
158
0.26
Class 1




420
830
8.14
169
0.26
Class 1




400
890
15.8
189
0.25
Class 1


Alloy 14
Ha
750
870
1.12
171
0.22
Class 1




710
910
2.38
180
0.13
Class 1




720
870
1.50
174
0.17
Class 1



Hb
620
850
4.45
209
0.14
Class 2



Hc
520
1340
10.76
143
0.79
Class 2




500
1290
10.10
166
0.80
Class 2




490
1220
9.15
159
0.70
Class 2



Hd
460
1310
11.30
140
0.98
Class 2




440
1310
12.00
184
0.97
Class 2




450
1320
12.54
154
0.94
Class 2



He
580
1230
8.54
155
0.67
Class 2




410
830
5.09
166
0.40
Class 2


Alloy 15
Ha
870
1080
1.51
203
N/A
Class 2




850
1180
2.98
186
0.21
Class 2




860
1130
1.94
173
0.23
Class 2



Hb
720
960
1.98
171
0.22
Class 1




730
920
1.59
183
0.22
Class 1



Hc
550
1090
10.23
184
0.54
Class 2




540
1140
10.94
191
0.56
Class 2




550
880
7.56
200
0.35
Class 2


Alloy 16
Hb
940
1290
2.01
168
0.26
Class 2



Hc
990
1260
1.57
178
N/A
Class 2




980
1270
1.77
183
N/A
Class 2


Alloy 17
He
500
1150
7.32
191
0.60
Class 2




500
1200
8.04
148
0.61
Class 2




480
1140
7.12
169
0.55
Class 2



Hc
490
1280
10.39
157
0.95
Class 2




430
1280
10.68
163
0.93
Class 2




480
1310
10.86
169
0.99
Class 2



Hd
440
1340
16.13
185
0.96
Class 2




430
1270
11.74
178
0.98
Class 2


Alloy 18
He
490
1280
8.70
148
0.73
Class 2




470
1000
5.80
154
0.55
Class 2



Hc
430
1230
9.66
223
0.70
Class 2




490
1290
10.81
160
0.99
Class 2




460
1300
11.29
156
0.95
Class 2



Hd
440
1270
16.70
154
0.89
Class 2




450
1240
12.39
139
0.99
Class 2




420
1270
13.51
157
0.95
Class 2


Alloy 19
He
550
1250
8.36
135
0.60
Class 2




570
1200
8.20
175
0.54
Class 2



Hc
480
1260
10.12
143
0.93
Class 2




510
1130
8.55
145
0.88
Class 2



Hd
460
1300
13.11
125
0.77
Class 2




490
1380
14.98
146
0.79
Class 2




440
1340
13.23
230
0.98
Class 2



Hf
430
1260
12.41
124
0.68
Class 2




440
1260
11.69
141
0.99
Class 2




390
1350
17.98
201
0.90
Class 2




440
1290
13.11
136
0.97
Class 2




430
1030
8.83
186
0.95
Class 2


Alloy 20
He
500
990
14.26
175
0.19
Class 1




490
950
12.42
170
0.20
Class 1




470
880
5.57
178
0.23
Class 1



Hc
470
990
17.66
171
0.21
Class 2




480
950
15.49
183
0.19
Class 2




480
950
15.69
169
0.20
Class 2



Hd
410
810
12.11
162
0.21
Class 2




430
920
16.83
155
0.22
Class 2


Alloy 21
He
440
910
5.82
186
0.26
Class 1




470
940
5.88
224
0.26
Class 1




470
880
5.07
168
0.28
Class 1



He
390
910
18.40
169
0.26
Class 1




440
920
10.96
176
0.25
Class 1




440
910
8.94
178
0.26
Class 1



Hd
380
890
19.38
192
0.26
Class 1




380
900
21.69
153
0.27
Class 1




360
910
24.41
145
0.27
Class 1


Alloy 22
He
650
1050
9.17
170
0.16
Class 2




620
1020
8.79
172
0.15
Class 2




600
1040
9.08
188
0.16
Class 2



Hc
540
1080
12.36
171
0.63
Class 2




540
980
11.05
163
0.41
Class 2




530
830
8.18
147
0.33
Class 2



Hd
480
1270
19.38
158
0.83
Class 2


Alloy 23
He
650
1390
3.37
179
0.45
Class 2




630
1430
3.84
175
0.46
Class 2



Hc
620
1250
2.59
140
0.51
Class 2




570
910
1.43
142
N/A
Class 2




690
1150
1.74
198
0.44
Class 2



Hd
550
1400
7.12
154
0.44
Class 2




630
1440
5.14
167
0.34
Class 2




660
1370
3.49
190
0.43
Class 2


Alloy 24
He
470
960
11.80
172
0.21
Class 1




510
860
3.91
206
0.25
Class 1




440
910
6.09
196
0.23
Class 1



Hc
450
920
15.94
174
0.20
Class 2




460
930
16.05
156
0.21
Class 2




450
990
19.24
148
0.22
Class 2



Hd
400
1010
23.05
165
0.26
Class 2




410
960
19.83
186
0.24
Class 2




440
1000
22.30
178
0.24
Class 2









Sheet Properties of HIPed and Heat Treated Sheets
After HIPing, the sheet material was heat treated in a box furnace at parameters specified in Table 8. The preferred aspect of the heat treatment after HIP cycle was to estimate thermal stability and property changes of the alloys by mimicking Stage 3 of the Twin Roll Casting process and also Stage 3 of the Thin Slab Casting process.







TABLE 8







Heat Treatment Parameters











Heat






Treatment

Temperature
Time



(ID)
Type
(° C.)
(min)
Cooling














T1
Age Hardening/Spinodal 
350
20
In air



Decomposition





T2
Age Hardening/Spinodal 
475
20
In air



Decomposition





T3
Age Hardening/Spinodal
600
20
In air



Decomposition





T4
Age Hardening/Spinodal 
700
20
In air



Decomposition





T5
Age Hardening/Spinodal
700
60
In air



Decomposition





T6
Age Hardening/Spinodal 
700
60
With 



Decomposition


furnace









The tensile specimens were cut from the sheets after HIP cycle and heat treatment using wire electrical discharge machining (EDM). Tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving; the load cell is attached to the top fixture. In Table 9, a summary of the tensile test results including tensile elongation, yield stress, ultimate tensile strength, Elastic Modulus and strain hardening exponent value are shown for the cast sheets after HIP cycle and heat treatment. As can be seen the tensile strength values vary from 530 to 1580 MPa. The tensile elongation varies from 0.71 to 30.24% and was observed to depend on alloy chemistry, HIP cycle, and heat treatment parameters which preferably determine microstructural formation in the sheets. Note that further increases in ductility up to 50% would be expected based on optimization of processing to eliminate further defects, especially casting defects which are present as pores in some of these sheets. Elastic Modulus was measured in a range from 104 to 267 GPa. Mechanical characteristic values strongly depend on alloy chemistry, HIP cycle parameters and heat treatment parameters. Strain hardening coefficient was calculated from the yield strength to the tensile strength resulting in ranges from 0.11 to 0.99 depending on alloy chemistry, structural formation, and different heat treatments.







TABLE 9







Summary on Tensile Test Results for Cast Sheets after HIP Cycle and Heat Treatment



















Ultimate







HIP
Heat
Yield
Tensile
Tensile
Elastic
Strain




Cycle
Treatment
Stress
Strength
Elongation
Modulus
Hardening
Type of


Alloy
ID
ID
(MPa)
(MPa)
(%)
(GPa)
Exponent
Behavior


















Alloy 1
Ha
T1
430
800
3.46
180
0.28
Class 1





430
850
4.81
184
0.27
Class 1




T2
440
790
2.60
200
0.29
Class 1





440
730
2.19
197
0.27
Class 1




T3
440
800
3.48
176
0.28
Class 1





410
870
7.14
165
0.28
Class 1



Hb
T1
430
720
3.45
182
0.26
Class 1





400
820
7.20
181
0.27
Class 1




T2
370
770
5.79
166
0.28
Class 1





410
860
8.25
187
0.26
Class 1




T3
390
830
7.36
174
0.28
Class 1





390
770
5.70
165
0.29
Class 1



Hc
T1
350
830
21.53
159
0.26
Class 1





340
810
21.35
148
0.26
Class 1





350
800
17.88
165
0.26
Class 1




T2
360
640
3.74
207
0.27
Class 1





390
840
17.59
129
0.25
Class 1




T3
340
800
21.63
143
0.27
Class 1





370
840
19.72
193
0.26
Class 1





360
680
5.45
198
0.27
Class 1


Alloy 2
Ha
T1
400
810
4.49
168
0.27
Class 1





400
840
6.10
153
0.28
Class 1




T2
400
740
3.30
207
0.29
Class 1





400
770
3.39
146
0.19
Class 1




T3
400
880
9.79
196
0.27
Class 1





400
660
2.57
146
0.29
Class 1





500
940
10.18
199
0.24
Class 1



Hb
T1
500
970
13.69
183
0.24
Class 1





500
890
8.50
162
0.26
Class 1





400
770
4.02
173
0.28
Class 1




T2
500
800
4.58
173
0.25
Class 1





500
940
10.32
133
0.25
Class 1





400
930
20.92
187
0.25
Class 1




T3
400
940
11.11
168
0.25
Class 1





500
810
4.96
118
0.28
Class 1





400
840
12.72
172
0.26
Class 1



Hc
T1
400
900
18.96
188
0.25
Class 1





400
680
4.96
151
0.29
Class 1





400
880
16.00
182
0.25
Class 1




T2
400
830
12.07
163
0.26
Class 1





400
860
11.52
198
0.25
Class 1




T3
400
900
19.25
185
0.26
Class 1





400
770
10.96
155
0.26
Class 1





400
850
18.48
168
0.26
Class 1


Alloy 3
Ha
T1
430
850
5.94
174
0.28
Class 1





420
860
7.01
165
0.27
Class 1





430
720
3.16
172
0.29
Class 1




T2
430
790
4.01
168
0.28
Class 1





420
790
4.08
173
0.28
Class 1





430
720
2.03
193
0.30
Class 1




T3
400
680
1.84
188
0.29
Class 1





400
850
4.96
174
0.30
Class 1





410
750
3.20
155
0.30
Class 1



Hb
T1
420
930
10.74
182
0.25
Class 1





420
930
12.71
182
0.25
Class 1





410
900
11.31
172
0.27
Class 1




T2
420
910
11.57
178
0.26
Class 1





410
920
12.26
183
0.26
Class 1





420
890
8.01
173
0.27
Class 1




T3
420
880
7.83
183
0.27
Class 1





400
890
8.52
196
0.27
Class 1





400
900
11.96
172
0.27
Class 1



Hc
T1
360
680
5.67
158
0.27
Class 1





370
690
4.27
169
0.28
Class 1





360
830
14.38
169
0.26
Class 1




T2
350
730
7.76
158
0.27
Class 1





360
820
19.95
167
0.25
Class 1




T3
360
530
2.68
176
0.28
Class 1





370
830
18.76
166
0.26
Class 1


Alloy 4
Ha
T1
600
820
1.21
183
N/A
Class 1





600
1020
3.26
180
0.28
Class 1





580
870
1.79
186
0.32
Class 1




T2
600
880
1.67
177
N/A
Class 1





620
830
1.11
197
N/A
Class 1





580
1040
3.32
182
0.29
Class 1




T3
620
1030
2.67
191
0.28
Class 1





600
1060
3.24
187
0.30
Class 1





590
980
3.44
164
0.29
Class 1



Hb
T1
530
940
2.84
170
0.31
Class 1





580
960
2.77
156
0.31
Class 1




T2
540
940
2.89
196
0.30
Class 1





570
1050
4.73
182
0.28
Class 1




T3
540
1030
4.74
175
0.29
Class 1





540
970
3.13
189
0.31
Class 1



Hc
T1
510
970
6.85
167
0.26
Class 1





490
930
5.29
196
0.27
Class 1





480
970
6.60
191
0.27
Class 1




T2
500
990
7.93
176
0.26
Class 1





490
950
6.36
173
0.27
Class 1




T3
490
970
8.16
187
0.26
Class 1





500
940
5.59
167
0.28
Class 1


Alloy 5
Hb
T1
500
850
2.81
168
0.30
Class 1





520
830
2.42
165
0.30
Class 1




T2
490
850
3.08
171
0.30
Class 1





540
850
2.31
166
0.29
Class 1




T3
500
880
3.52
171
0.29
Class 1



Hc
T1
450
710
2.29
186
0.29
Class 1





490
950
7.98
186
0.25
Class 1





470
880
5.75
199
0.26
Class 1




T2
460
940
7.65
197
0.26
Class 1





470
970
11.06
170
0.25
Class 1





460
950
9.12
190
0.26
Class 1




T3
480
950
8.95
191
0.25
Class 1





460
960
10.44
180
0.25
Class 1


Alloy 6
Ha
T1
550
880
2.15
194
0.29
Class 1




T2
570
940
2.63
185
0.29
Class 1




T3
540
910
2.69
205
0.28
Class 1





600
980
2.66
203
0.28
Class 1



Hb
T1
540
790
1.54
194
N/A
Class 1





560
920
2.45
198
0.28
Class 1





500
800
1.78
183
0.31
Class 1




T2
550
790
1.44
180
N/A
Class 1





530
880
2.38
170
0.30
Class 1





540
820
1.97
191
0.29
Class 1




T3
520
970
3.87
186
0.28
Class 1





550
970
3.24
180
0.30
Class 1



Hc
T1
460
950
8.93
199
0.25
Class 1





480
950
7.21
173
0.26
Class 1




T2
490
970
8.62
180
0.25
Class 1





480
960
7.20
186
0.26
Class 1





480
940
6.98
177
0.27
Class 1




T3
460
940
9.55
193
0.25
Class 1





460
960
7.55
172
0.26
Class 1





470
980
8.63
170
0.26
Class 1


Alloy 7
Ha
T1
570
950
2.46
191
0.30
Class 1





570
770
1.21
178
N/A
Class 1




T2
620
900
2.13
188
0.26
Class 1





570
910
2.04
203
0.29
Class 1




T3
580
930
2.35
187
0.30
Class 1





590
960
2.55
192
0.28
Class 1



Hb
T1
560
990
3.36
167
0.30
Class 1





520
720
1.24
175
N/A
Class 1




T2
510
830
1.83
177
0.33
Class 1





500
840
2.58
136
0.34
Class 1





520
840
2.07
213
0.30
Class 1




T3
540
850
1.84
195
0.31
Class 1



Hc
T1
480
800
2.38
202
0.29
Class 1





480
950
6.07
167
0.27
Class 1




T2
500
820
2.38
209
0.29
Class 1





450
680
1.60
158
N/A
Class 1




T3
480
840
3.01
152
0.32
Class 1





500
930
5.16
156
0.28
Class 1


Alloy 8
Ha
T1
580
950
2.17
229
0.30
Class 1




T2
620
910
1.61
186
N/A
Class 1





640
1030
2.53
172
0.30
Class 1




T3
650
930
1.68
185
N/A
Class 1



Hb
T1
580
1030
3.27
183
0.30
Class 1





590
1040
4.10
149
0.30
Class 1




T2
560
970
3.20
151
0.31
Class 1





560
980
2.77
181
0.31
Class 1





580
850
1.72
172
0.32
Class 1




T3
540
910
2.16
166
0.33
Class 1





580
1040
3.59
201
0.29
Class 1



Hc
T1
500
950
4.55
186
0.28
Class 1





510
810
2.04
181
0.31
Class 1




T2
500
770
1.87
169
0.31
Class 1





520
990
6.06
177
0.28
Class 1




T3
470
580
0.90
138
N/A
Class 1





510
1000
7.32
162
0.27
Class 1





350
560
1.07
213
N/A
Class 1


Alloy 10
Hb
T1
550
960
3.09
170
0.32
Class 1





530
800
1.76
176
0.32
Class 1




T2
510
1040
5.16
161
0.31
Class 1





540
720
1.32
183
0.31
Class 1




T3
530
850
2.23
171
0.32
Class 1


Alloy 11
Hb
T1
500
1180
6.85
170
0.87
Class 2





480
920
4.94
172
0.50
Class 2




T2
490
1040
6.18
166
0.88
Class 2





460
900
4.75
179
0.66
Class 2




T3
470
1050
5.81
182
0.87
Class 2





430
1050
5.21
160
0.81
Class 2



Hc
T1
700
1290
5.84
161
0.34
Class 2





880
1360
5.24
186
0.25
Class 2





840
1390
7.44
187
0.28
Class 2




T2
480
1070
5.12
170
0.52
Class 2





990
1140
2.44
166
N/A
Class 2





860
1410
6.66
163
0.40
Class 2




T3
530
1260
8.65
169
0.49
Class 2





400
1190
5.40
169
0.92
Class 2





430
1070
3.49
159
0.67
Class 2


Alloy 12
Hb
T1
460
880
4.58
161
0.28
Class 1





420
780
3.71
181
0.28
Class 1




T2
430
780
3.48
169
0.30
Class 1





440
820
4.49
163
0.28
Class 1




T3
420
740
2.75
193
0.30
Class 1





400
830
4.17
185
0.28
Class 1



Hc
T1
380
850
10.45
177
0.26
Class 1





370
880
16.32
185
0.25
Class 1




T2
420
870
10.49
146
0.25
Class 1





400
850
8.48
176
0.26
Class 1




T3
400
850
10.38
168
0.26
Class 1





390
850
10.28
159
0.25
Class 1


Alloy 13
Hb
T1
470
800
2.98
168
0.29
Class 1





490
560
1.33
181
N/A
Class 1




T2
430
780
4.09
176
0.27
Class 1




T3
430
620
1.74
183
N/A
Class 1





470
800
2.98
168
0.29
Class 1



Hc
T1
400
890
15.28
168
0.25
Class 1





420
880
12.08
158
0.25
Class 1




T2
410
860
11.06
170
0.26
Class 1





410
840
10.23
187
0.25
Class 1




T3
400
860
12.88
155
0.26
Class 1





410
880
12.70
148
0.26
Class 1





400
890
16.48
163
0.25
Class 1


Alloy 14
Ha
T1
730
840
1.39
157
N/A
Class 1





700
940
4.32
172
0.11
Class 1





740
980
4.73
168
0.11
Class 1




T2
690
820
1.07
186
N/A
Class 1





710
910
2.57
167
0.13
Class 1




T3
680
810
1.61
153
N/A
Class 1





670
850
2.68
154
0.15
Class 1



Hb
T1
630
1040
6.77
163
0.47
Class 2





620
1010
6.42
178
0.46
Class 2




T2
640
980
6.04
158
0.41
Class 2





640
1120
7.54
151
0.57
Class 2




T3
600
690
1.22
182
0.54
Class 2





650
1090
7.00
156
0.54
Class 2





620
1070
6.78
171
0.56
Class 2



Hc
T1
520
1150
8.28
164
0.66
Class 2





520
1350
11.00
179
0.88
Class 2





500
1190
8.75
134
0.87
Class 2




T2
520
1320
10.04
191
0.77
Class 2





470
1170
8.49
169
0.88
Class 2




T3
490
1350
10.24
122
0.82
Class 2





490
1160
7.96
170
0.93
Class 2





500
1400
12.67
174
0.87
Class 2



Hd
T1
420
1250
12.52
129
0.99
Class 2





440
1320
12.87
159
0.93
Class 2





410
910
7.73
128
0.81
Class 2




T2
370
930
8.07
148
0.88
Class 2





420
1050
8.66
126
0.91
Class 2




T3
430
1320
13.55
129
0.94
Class 2





440
1300
12.30
139
0.98
Class 2





440
830
6.59
186
0.80
Class 2




T4
400
1160
9.22
92
0.97
Class 2





400
1280
11.15
137
0.95
Class 2





380
1330
12.98
123
0.95
Class 2




T5
410
1300
10.35
140
0.97
Class 2




T6
410
1320
11.23
167
0.93
Class 2





380
1310
13.50
160
0.91
Class 2



He
T1
560
1100
7.37
164
0.59
Class 2





590
1040
6.66
159
0.53
Class 2




T2
560
1140
7.70
159
0.61
Class 2





560
960
5.96
169
0.50
Class 2




T3
530
1050
6.60
167
0.60
Class 2





550
1070
6.80
148
0.63
Class 2


Alloy 15
Hc
T1
600
1100
10.15
158
0.64
Class 2





560
950
8.66
187
0.46
Class 2




T2
600
1040
9.68
176
0.56
Class 2





550
1000
9.23
174
0.53
Class 2




T3
360
1120
10.73
146
0.71
Class 2





560
940
8.27
189
0.54
Class 2


Alloy 16
Hb
T1
1130
1570
4.18
235
0.19
Class 2




T2
960
1160
0.71
222
N/A
Class 2





1280
1580
2.41
193
0.21
Class 2




T3
1070
1200
1.65
202
0.15
Class 2





1130
1300
1.71
220
0.16
Class 2





1140
1420
6.06
209
0.13
Class 2



Hc
T1
1070
1270
1.26
175
N/A
Class 2





990
1160
0.70
203
N/A
Class 2





750
1420
2.42
183
0.21
Class 2




T2
1110
1210
0.74
198
N/A
Class 2





1290
1500
1.58
ISO
0.24
Class 2





1070
1260
0.86
328
0.30
Class 2




T3
980
1170
2.79
189
0.14
Class 2





1080
1260
4.14
222
0.10
Class 2





1080
1200
2.04
190
0.12
Class 2


Alloy 17
He
T4
550
1300
9.21
166
0.76
Class 2





550
1280
8.89
184
0.77
Class 2





510
1210
7.80
142
0.69
Class 2




T5
530
1310
9.80
154
0.73
Class 2





540
1230
7.98
176
0.80
Class 2





470
1200
7.89
176
0.68
Class 2




T6
550
1170
7.72
125
0.52
Class 2





490
1200
7.69
170
0.54
Class 2





510
1350
10.27
127
0.62
Class 2



Hd
T4
430
1320
13.06
186
0.97
Class 2





440
1310
13.81
157
0.92
Class 2





420
1280
10.20
165
0.93
Class 2




T5
400
1300
16.03
116
0.92
Class 2





390
1300
13.44
182
0.98
Class 2





400
1300
12.58
169
0.99
Class 2




T6
400
1290
11.11
132
0.98
Class 2





400
1300
12.21
160
0.89
Class 2



Hc
T4
490
1260
9.74
ISO
0.87
Class 2





480
1360
12.92
176
0.90
Class 2





490
1300
10.75
148
0.78
Class 2




T5
430
1170
9.07
121
0.79
Class 2





470
1340
11.37
128
0.83
Class 2





460
1360
12.03
164
0.98
Class 2




T6
450
1360
12.07
170
0.97
Class 2





470
1290
10.06
157
0.99
Class 2





440
1290
11.53
135
0.79
Class 2


Alloy 18
He
T4
470
1340
9.49
150
0.72
Class 2





500
1290
8.55
151
0.74
Class 2





490
1380
11.44
146
0.73
Class 2




T5
450
1360
10.41
162
0.66
Class 2





440
1290
8.51
161
0.64
Class 2





440
1330
9.71
159
0.67
Class 2




T6
480
1240
7.49
180
0.67
Class 2





420
1350
10.16
194
0.68
Class 2





480
1320
9.60
114
0.69
Class 2



Hc
T4
450
1270
10.40
185
0.98
Class 2





460
1320
11.56
172
0.99
Class 2




T5
430
1250
9.00
177
0.90
Class 2





450
1290
9.57
182
0.99
Class 2




T6
430
1310
15.40
152
0.84
Class 2





420
1330
16.03
147
0.88
Class 2



Hd
T4
420
1170
9.99
144
0.98
Class 2





440
1290
16.05
104
0.91
Class 2





370
1240
11.34
163
0.98
Class 2




T5
380
1290
14.91
131
0.86
Class 2





400
1290
12.67
118
0.86
Class 2





400
1290
14.93
136
0.89
Class 2




T6
380
1260
12.01
120
0.86
Class 2





360
1300
18.80
112
0.83
Class 2





360
1270
11.15
146
0.86
Class 2


Alloy 19
He
T4
570
1200
7.80
162
0.68
Class 2





590
1260
8.18
154
0.71
Class 2





580
1290
8.49
175
0.67
Class 2




T5
560
1270
8.23
139
0.68
Class 2





550
1070
6.68
188
0.65
Class 2





570
950
5.80
172
0.50
Class 2




T6
540
1310
9.16
150
0.77
Class 2





560
1100
6.82
170
0.63
Class 2



Hc
T4
480
1160
8.44
138
0.86
Class 2





530
1160
8.35
143
0.79
Class 2




T5
480
1300
8.72
172
0.98
Class 2





390
900
6.03
154
0.72
Class 2




T6
450
1030
6.18
169
0.56
Class 2





470
1270
7.93
150
0.71
Class 2





380
940
5.83
160
0.50
Class 2



Hd
T4
480
1390
18.51
141
0.84
Class 2





460
1380
18.19
174
0.87
Class 2





500
1380
14.89
116
0.89
Class 2




T5
450
1370
16.27
180
0.88
Class 2





470
1330
10.96
205
0.97
Class 2





400
1370
17.69
195
0.91
Class 2




T6
430
1370
16.60
122
0.81
Class 2





430
1360
15.02
139
0.81
Class 2





450
1350
14.64
150
0.83
Class 2



Hf
T4
430
1360
18.66
145
0.91
Class 2





430
1220
13.4
267
N/A
Class 2





380
1350
14.75
256
0.95
Class 2




T5
400
1350
15.29
153
0.97
Class 2





360
1350
14.19
171
0.98
Class 2





390
1240
9.48
143
0.80
Class 2




T6
370
1340
18.48
136
0.82
Class 2





390
1340
13.95
128
0.90
Class 2





360
1330
17.02
135
0.79
Class 2


Alloy 20
He
T4
490
920
6.94
169
0.20
Class 1





520
1050
17.47
179
0.19
Class 1





490
1010
16.92
181
0.19
Class 1




T5
500
970
12.71
185
0.17
Class 2





540
980
13.52
168
0.19
Class 2





500
910
7.49
171
0.21
Class 2




T6
460
860
4.72
154
0.26
Class 2





500
990
14.58
129
0.19
Class 2





530
990
13.22
155
0.19
Class 2



Hc
T4
470
960
15.19
156
0.19
Class 2





410
1090
22.28
176
0.27
Class 2





440
970
16.18
167
0.20
Class 2




T5
470
950
15.12
178
0.20
Class 2





460
910
13.33
180
0.17
Class 2





470
960
14.78
165
0.19
Class 2




T6
460
880
12.17
166
0.17
Class 2





500
1060
18.71
198
0.25
Class 2





500
1070
17.52
174
0.26
Class 2



Hd
T4
440
950
17.41
167
0.23
Class 2





450
920
16.55
181
0.22
Class 2





470
990
20.19
138
0.28
Class 2




T5
420
1050
22.42
179
0.31
Class 2





440
1020
22.04
179
0.31
Class 2




T6
420
950
19.50
168
0.27
Class 2





440
1010
20.63
174
0.30
Class 2


Alloy 21
He
T4
420
960
8.18
182
0.25
Class 1





500
990
8.99
215
0.24
Class 1




T5
460
900
5.94
195
0.26
Class 1





470
970
8.64
248
0.24
Class 1





490
960
7.79
165
0.26
Class 1




T6
410
1000
10.11
221
0.25
Class 1





460
980
10.63
186
0.25
Class 1





510
990
8.73
141
0.26
Class 1



Hc
T4
430
970
15.00
184
0.23
Class 1





410
880
9.42
172
0.24
Class 1





430
910
9.18
159
0.25
Class 1




T5
430
930
13.58
170
0.25
Class 1





430
950
13.24
170
0.24
Class 1





430
920
10.24
162
0.26
Class 1




T6
430
880
7.08
177
0.27
Class 1





430
960
14.89
171
0.25
Class 1





430
970
17.95
184
0.25
Class 1



Hd
T4
400
920
26.12
185
0.25
Class 1





380
910
24.16
156
0.26
Class 1




T5
390
940
30.24
165
0.26
Class 1





410
930
21.97
126
0.25
Class 1





390
930
27.70
140
0.25
Class 1




T6
360
860
14.74
179
0.25
Class 1





370
910
19.52
157
0.26
Class 1





390
930
25.58
181
0.25
Class 1


Alloy 22
He
T4
610
910
6.11
204
0.11
Class 2





630
1100
9.88
156
0.19
Class 2





650
930
7.05
187
0.12
Class 2




T5
670
1100
10.01
165
0.37
Class 2





420
980
7.55
221
0.22
Class 2





590
1020
8.33
189
0.27
Class 2




T6
660
860
3.86
149
0.13
Class 2





620
980
8.15
121
0.16
Class 2





650
1170
10.95
169
0.20
Class 2



Hc
T4
550
1260
15.93
160
0.68
Class 2





530
1260
15.88
163
0.68
Class 2




T5
530
1250
14.60
168
0.76
Class 2





530
970
10.06
165
0.55
Class 2




T6
520
1180
14.95
132
0.60
Class 2





580
1320
18.91
120
0.71
Class 2





510
840
7.91
189
0.16
Class 2



Hd
T4
480
1270
19.77
140
0.80
Class 2





470
1120
14.22
154
0.74
Class 2





500
1270
19.73
118
0.81
Class 2




T5
410
930
10.57
176
0.82
Class 2





430
1010
11.95
177
0.79
Class 2





480
1140
13.78
130
0.79
Class 2




T6
480
1260
19.48
143
0.80
Class 2





460
880
10.01
154
0.47
Class 2





490
1210
16.19
155
0.76
Class 2


Alloy 23
He
T4
510
1100
3.90
240
0.45
Class 2





530
1170
4.36
183
0.50
Class 2




T5
670
1320
6.29
173
0.43
Class 2





680
1120
4.58
165
0.23
Class 2





620
1010
3.66
242
0.25
Class 2




T6
620
1100
2.18
172
0.46
Class 2





650
1390
4.57
142
0.41
Class 2





630
1250
3.11
146
0.47
Class 2



Hc
T4
500
960
3.24
166
0.46
Class 2




T6
730
1090
4.68
138
0.30
Class 2





630
1190
5.72
157
0.41
Class 2



Hd
T4
570
1370
9.54
126
0.45
Class 2





490
1360
8.53
153
0.53
Class 2





540
1250
4.25
159
0.43
Class 2




T5
640
1350
9.19
177
0.30
Class 2





610
1350
7.96
191
0.29
Class 2




T6
660
1300
12.64
136
0.40
Class 2





690
1300
7.86
167
0.40
Class 2





670
1340
12.10
179
0.40
Class 2


Alloy 24
He
T4
450
930
10.52
169
0.16
Class 1





470
930
8.27
181
0.22
Class 1





500
930
9.54
192
0.20
Class 1




T5
410
880
5.23
245
0.23
Class 1





510
930
9.90
195
0.19
Class 1





500
910
10.45
148
0.20
Class 1




T6
490
810
2.68
184
0.26
Class 1





490
810
3.88
170
0.23
Class 1





560
960
9.43
143
0.12
Class 1



Hc
T4
470
1050
20.86
170
0.23
Class 2





440
910
15.19
177
0.20
Class 2





460
830
9.10
178
0.21
Class 2




T5
460
930
15.09
164
0.21
Class 2





370
910
15.18
130
0.23
Class 2





450
650
2.11
199
0.25
Class 2




T6
460
950
15.59
171
0.20
Class 2





460
1080
22.31
173
0.29
Class 2



Hd
T4
410
900
17.13
158
0.24
Class 2





410
1070
26.26
152
0.29
Class 2





410
980
20.70
156
0.26
Class 2




T5
400
790
12.61
172
0.19
Class 2





410
1080
26.25
157
0.38
Class 2





410
1040
21.27
163
0.32
Class 2




T6
410
1040
22.79
146
0.33
Class 2





400
810
11.94
160
0.20
Class 2





410
1020
21.28
163
0.32
Class 2









Comparative Examples
Case Example #1
Tensile Properties Comparison with Existing Steel Grades
Tensile properties of selected alloy were compared with tensile properties of existing steel grades. The selected alloys and corresponding treatment parameters are listed in Table 10. Tensile stress—strain curves are compared to that of existing Dual Phase (DP) steels (FIG. 7); Complex Phase (CP) steels (FIG. 8); Transformation Induced Plasticity (TRIP) steels (FIG. 9); and Martensitic (MS) steels (FIG. 10). A Dual Phase Steel may be understood as a steel type consisting of a ferritic matrix containing hard martensitic second phases in the form of islands, a Complex Phase Steel may be understood as a steel type consisting of a matrix consisting of ferrite and bainite containing small amounts of martensite, retained austenite, and pearlite, a Transformation Induced Plasticity steel may be understood as a steel type which consists of austenite embedded in a ferrite matrix which additionally contains hard bainitic and martensitic second phases and a Martensitic steel may be understood as a steel type consisting of a martensitic matrix which may contain small amounts of ferrite and/or bainite. As it can be seen, the alloys claimed in this disclosure have superior properties as compared to existing advanced high strength (AHSS) steel grades.







TABLE 10







Downselected Tensile Curves Labels and Identity












Curve






Label
Alloy
HIP
HT







A
Alloy 16
 850° C. for 1 hour
350° C. for 20 min



B
Alloy 23
1100° C. for 1 hour
None



C
Alloy 14
1000° C. for 1 hour
650° C. for 20 min



D
Alloy 19
1100° C. for 1 hour
700° C. for 20 min



E
Alloy 22
1100° C. for 1 hour
700° C. for 20 min



F
Alloy 24
1100° C. for 1 hour
700° C. for 20 min



G
Alloy 21
1100° C. for 1 hour
700° C. for 1 hr










Case Example #2
Modal Structure
Microstructure of the sheets from selected alloys with chemical composition specified in Table 2 in as-cast state, after HIP cycle and after HIP cycle with additional heat treatment was examined by scanning electron microscopy (SEM) using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. Examples of Modal Structure (Structure #1) and NanoModal Structures (Structure #2) in selected alloys are shown in FIGS. 11 through 15. As it can be seen, the Modal structure may be formed in alloys in as-cast state (FIG. 11). To produce the NanoModal Structure additional thermal mechanical treatment might be needed such as HIP cycle (FIGS. 12-13) and/or HIP cycle with additional heat treatment (FIGS. 14 and 15). Other types of thermal mechanical treatment such as hot rolling, forging, hot stamping, etc., might be also effective for NanoModal Structure formation in the alloys with referenced chemistries described in this application. Formation of modal structure in sheet materials is the first step in achieving high ductility at moderate strength (Class 1 steels) while achieving the NanoModal Structure is enabling for Class 2 steels.
Case Example #3
Structure Development in Alloy 1
According to the alloy stoichiometries in Table 2, the Alloy 1 was weighed out from high purity elemental charges. It should be noted that Alloy 1 has demonstrated Class I behavior with high plastic ductility at moderate strength. The resulting charges were arc-melted into 4 thirty-five gram ingots and flipped and re-melted several times to ensure homogeneity. The resulting ingots were then re-melted and cast into 3 sheets under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. An example picture of one of the 1.8 mm thick Alloy 1 sheets is shown in FIG. 16. Two of the sheets were then HIPed at 1000° C. for 1 hour. One of the HIPed sheets was then subsequently heat treated at 350° C. for 20 minutes. The sheets including as-cast, HIPed and HIPed/heat treated ones were then cut up using a wire-EDM to produce samples for various studies including tensile testing, SEM microscopy, TEM microscopy, and X-ray diffraction.
Samples that were cut out of the Alloy 1 sheets were metallographically polished in stages down to 0.02 μm Grit to ensure smooth samples for scanning electron microscopy (SEM) analysis. SEM was done using a Zeiss EVO-MA10 model with the maximum operating voltage of 30 kV. Example SEM backscattered electron micrographs of the Alloy 1 sheet samples in the as-cast, HIPed and HIPed and heat treated conditions are shown in FIG. 17.
As shown, the microstructure of the Alloy 1 sheet exhibits Modal Structures in all three conditions. In the as-cast sample, three areas can be readily identified (FIG. 17a). The matrix phase in a form of individual grains of 5 to  ̃10 μm in size are marked by #3 in FIG. 17a. These grains are separated by intergranular regions (#2 in FIG. 17a). Additional isolated precipitates are marked by #1 in FIG. 17a. The black phase precipitates (#1) represent a high Si-containing phase as identified by energy-dispersive spectroscopy (EDS). The intergranular region (#2) apparently contains higher concentration of light elements (such as B, Si) as compared to matrix grains #3. After the HIP cycle, significant change occurs in the intergranular region (#2). A number of fine precipitates, which are typically less than 500 nm in size, form in this area (FIG. 17b). These precipitates are predominantly distributed in the intergranular region #2, while matrix grains #3 and precipitates #1 do not show obvious change in terms of morphology and size. After heat treatment, the microstructure appears to be similar to that after HIP cycle, but additional finer precipitates are formed (FIG. 17c).
Additional details of the Alloy 1 sheet structure are revealed by using X-ray diffraction. X-ray diffraction was done using a Panalytical X′Pert MPD diffractometer with a Cu Kα X-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. The resulting scans were then subsequently analyzed using Rietveld analysis using Siroquant software. In FIGS. 18-20, X-ray diffraction scan patterns are shown including the measured/experimental pattern and the Rietveld refined pattern for the Alloy 1 sheets in the as-cast, HIPed, and HIPed/heat treated conditions, respectively. As can be seen, good fits of the experimental data were obtained in all cases. Analysis of the X-ray patterns including specific phases found, their space groups and lattice parameters are shown in Table 11. Note that the space group represents a description of the symmetry of the crystal and can have one of 230 types and is further identified with its corresponding Hermann Maugin space group symbol. In all cases, two phases were found, a cubic γ-Fe (austenite) and a complex mixed transitional metal boride phase with the M2B stoichiometry. Note that while a third phase appears to exist from the SEM microscopy studies, this phase was not identified by the X-ray diffraction scans indicating that intergranular region might be represented by a fine mixture of two identified phases. Note also that the lattice parameters of the identified phases are different than that found for pure phases clearly indicating the effects of dissolution by the alloying elements. For example, γ-Fe as a pure phase would exhibit a lattice parameter equal to a=3.575 Å and Fe2B pure phase would exhibit lattice parameters equal to a=5.099 Å and c=4.240 Å. Note that based on the significant change in lattice parameters in the M2B phase is it likely that silicon is also dissolved into this structure so it is not a pure boride phase. Additionally, as can be seen in Table 11, while the phases do not change, the lattice parameters do change as a function of the condition of the sheet (i.e. cast, HIPed, HIPed and heat treated) which indicates that redistribution of alloying elements is occurring.
To examine the structural details of the Alloy 1 sheets in more detail, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM specimens, samples were cut from the as-cast, HIPed, and HIPed/heat-treated sheets. The samples were then ground and polished to a thickness of 30 ̃40 μm. Discs of 3 mm in diameter were punched from these thin samples, and the final thinning was done by twin-jet electropolishing using a 30% HNO3 in methanol base. The prepared specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope (TEM) operated at 200 kV.
In FIG. 21, TEM micrographs of the Alloy 1 sheet samples are shown for a) As-Cast, b) HIPed at 1000° C. for 1 hour, and c) HIPed at 1000° C. for 1 hour with subsequent heat treatment at 350° C. for 20 minutes, respectively. In the as-cast sample, the matrix grains are in the range of 5 ̃10 μm in size (FIG. 21a) that are consistent with the SEM observation in FIG. 17a. In addition, lamella structure is revealed in the intergranular regions that separate the matrix grains. The lamella structure corresponds to the area #2 in FIG. 17a. The lamella spacing is typically of  ̃200 nm, which is beyond the limit of SEM resolution and not seen in FIG. 17a. After HIP cycle, the lamella structure is re-organized into the isolated precipitates of less than 500 nm in size distributed in the region between matrix grains which retain the same size as in the as-cast sample (FIG. 21b). Unlike the lamellas, the precipitates are discontinuous indicating that significant microstructural changes were induced by HIP cycle. Heat treatment does not induce large changes in the microstructure, but some finer precipitates can be identified by TEM (FIG. 21c). As noted above, Alloy 1 behaves herein as a Class 1 Steel and there is no Static Nanophase Refinement or Dynamic Nanophase Strengthening observed.







TABLE 11







Rietveld Phase Analysis of Alloy 1 Sheet











Condition
Phase 1
Phase 2







As-Cast Sheet
γ-Fe
M2B




Structure: Cubic
Structure:




Space group #:
Tetragonal




#225
Space group #:




Space group:
#140




Fm3m
Space group:




LP: a = 3.588 Å
I4/mcm





LP: a = 5.168 Å





c = 4.201 Å



HIPed at 1000° C. 
γ-Fe
M2B



for 1 hour
Structure: Cubic
Structure:




Space group #:
Tetragonal




#225
Space group #:




Space group:
#140




Fm3m
Space group:




LP: a = 3.585 Å
I4/mcm





LP: a = 5.295 Å





c = 4.186 Å



HIPed at 1000° C. for 
γ-Fe
M2B



1 hour, Heat treated 
Structure: Cubic
Structure:



at 350° C. for 20 minutes
Space group #:
Tetragonal




#225
Space group #:




Space group:
#140




Fm3m
Space group:




LP: a = 3.585 Å
I4/mcm





LP: a = 5.177 Å





c = 4.234 Å










Case Example #4
Tensile Properties and Structural Changes in Alloy 1
The tensile properties of the steel sheet produced in this application will be sensitive to the specific structure and specific processing conditions that the sheet experiences. In FIG. 22, the tensile properties of Alloy 1 sheet representative of a Class 1 steel are shown in the as-cast, HIPed (1000° C. for 1 hour) and HIPed (1000° C. for 1 hour)/heat treated (350° C. for 20 minutes) conditions. As can be seen, the as-cast sheet shows relatively lower ductility than the HIPed and HIPed/heat treated samples. This increase in ductility may be attributed to both the reduction of macrodefects in the HIPed sheets and microstructural changes occurring in the modal structures of the HIPed or HIPed/heat treated sheets as discussed earlier in Case Example #3. Additionally, during the application of a stress during tensile testing, it will be shown that structural changes are occurring.
For the Alloy 1 sheet HIPed at 1000° C. for 1 hour and heat treated at 350° C. for 20 minutes, structural details were obtained through using X-ray diffraction which was done on both the undeformed sheet samples and on the gage sections of the deformed tensile specimens. X-ray diffraction was specifically done using a Panalytical X′Pert MPD diffractometer with a Cu Kα x-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. In FIG. 23, X-ray diffraction patterns are shown for the Alloy 1 sheet HIPed at 1000° C. for 1 hour and heat treated at 350° C. for 20 minutes in both the undeformed sheet and the gage section of the tensile tested sample cut out from the sheet. As can be readily seen, there are significant structural changes occurring during deformation with new phases formation as indicated by new peaks in the X-ray pattern. Peak shifts indicate that redistribution of alloying elements is occurring between the phases present in both samples.
The X-ray pattern for the deformed Alloy 1 tensile tested specimen (HIPed (1000° C. for 1 hour)/heat treated at 350° C. for 20 minutes) was subsequently analyzed using Rietveld analysis using Siroquant software. As shown in FIG. 24, a close agreement was found between the measured and calculated patterns. In Table 12, the phases identified in the Alloy 1 sheet before and after tensile deformation are compared. As can be seen, the γ-Fe and M2B phases are present in the sheet before and after tensile testing although the lattice parameters changed indicating that the amount of solute elements dissolved in this phases changed. Furthermore, as shown in Table 12, after deformation, two new previously unknown hexagonal phases have been identified. One newly identified hexagonal phase is representative of a dihexagonal pyramidal class and has a hexagonal P63mc space group (#186) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 25a. The other hexagonal phase is representative of a ditrigonal dipyramidal class and has a hexagonal P6bar2C space group (#190) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 25b. It is theorized based on the small crystal unit cell size that this phase is likely a silicon based phase possibly a previously unknown Si-B phase. Note that in the FIG. 25, key lattice planes are identified corresponding to significant Bragg diffraction peaks.







TABLE 12







Rietveld Phase Analysis of Alloy 1 Sheet; 


Before and After Tensile Testing











Condition
Phase 1
Phase 2
Phase 3
Phase 4





Shed - HIPed
γ-Fe
M2B




at 1000° C. for
Structure:
Structure:




1 hour and
Cubic
Tetragonal




heat treated at
Space 
Space 




350° C. for 20
group #:
group #:




minutes - Prior 
#225
#140




to tensile
Space group:
Space group;




testing
Fm3m
I4/mcm





LP: 
LP: 





a = 3.585 Å
a = 5.177 Å






c = 4.234 Å




Sheet -HIPed
γ-Fe
M2B
Hexagonal
Hexagonal


at 1000° C. for
Structure:
Structure:
Phase 1
Phase 2


1 hour and
Cubic
Tetragonal
(new)
(new)


heat treated at
Space 
Space 
Structure:
Structure:


350° C. for 20
group #:
group #:
Hexagonal
Hexagonal


minutes - After
#225
#140
Space 
Space 


tensile testing
Space group:
Space group:
group #:
group #:



Fm3m
I4/mcm
#186
#190



LP: 
LP: 
Space group:
Space group:



a = 3.589 Å
a = 5.290 Å
P63mc
P62barC




c = 4.204 Å
LP: 
LP: 





a = 2.870 Å
a = 4.995 Å





c = 6.079 Å
c = 11.374 Å









To focus on structural changes occurring during tensile testing, the Alloy 1 sheet HIPed at 1000° C. for 1 hour, and heat treated at 350° C. for 20 minutes was examined before and after deformation. TEM specimens were prepared from the undeformed HIPed and heat treated sheet and from the gage section of the sample cut off the same sheet and tested in tension until failure. TEM specimens were made from the sheet first by mechanical grinding/polishing, and then electrochemical polishing. TEM specimens of deformed tensile specimens were cut directly from the gage section and then prepared in an analogous manner to the undeformed sheet specimens. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV.
In FIG. 26, TEM micrographs of microstructure in undeformed sheet and in a gage section after the tensile testing are shown. In the undeformed sample, the matrix grains are very clean, free of defects such as dislocations due to the high temperature exposure during HIP cycle, but the precipitates in the intergranular region are clearly seen (FIG. 26a). After the tensile testing, a high density of dislocations was observed in the matrix grains. A number of dislocations were also pinned by the precipitates in the intergranular region. Additionally, some very fine precipitates appear (i.e. Dynamic Nanophase Formation) within the matrix grains after the tensile testing, as shown in FIG. 26b. These very fine precipitates may correspond to the new hexagonal and face centered cubic type phases identified by X-ray diffraction (see subsequent section). The new hexagonal phase could also form as fine precipitates in the intergranular region where an extensive deformation may also take place. Due to the pinning effect by the precipitates, the matrix grains do not change their geometry during the tensile deformation. While the deformation-induced nanoscale phase formation may contribute to the hardening in the Alloy 1 sheet, the work-hardening of Alloy 1 appears to be dominated by dislocation based mechanisms including dislocation pinning by precipitates.
The more detailed microstructure of the Alloy 1 sheet sample that was HIPed at 1000° C. for 1 hour, heat treated at 350° C. for 20 minutes, and, then tensile tested is shown in FIGS. 27-28. In the matrix grains, the dislocations of high density interact with each other forming dislocation cells. Occasionally, stacking faults and twins can be found in the grains as well. Meanwhile, the precipitates in the intergranular regions also pin down the dislocations, as shown in FIG. 27. Both in the grains and in the intergranular region, some very fine precipitates can be seen to form during the tensile deformation.
Due to micron sized matrix grains in the Alloy 1 sheet, the deformation is dominated by dislocation mechanism with corresponding strain hardening behavior. Some additional strain hardening may occur due to twining/stacking faults. A hexagonal phase formation corresponding to Dynamic Nanophase Strengthening (Mechanism #2) is also detected in the Alloy 1 sheet during the deformation. The Alloy 1 sheet is an example of Class 1 steel with Modal Structure formation and Dynamic Nanophase Strengthening leading to high ductility at moderate strength.
Case Example #5
Structure Development in Alloy 14
According to the alloy stoichiometries in Table 2, the Alloy 14 was weighed out using high purity elemental charges. I should be noted that Alloy 14 has demonstrated Class 2 behavior with high plastic ductility at high strength. The resulting charges were arc-melted into 4 thirty-five gram ingots and flipped and re-melted several times to ensure homogeneity. The resulting ingots were then re-melted and cast into 3 sheets under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. An example picture of one of the 1.8 mm thick Alloy 14 sheets is shown in FIG. 29. Two of the sheets were then HIPed at 1000° C. for 1 hour. One of the HIPed sheets was then subsequently heat treated at 350° C. for 20 minutes. The sheets in the as-cast, HIPed and HIPed/heat treated states were then cut up using a wire-EDM to produce samples for various studies including tensile testing, SEM microscopy, TEM microscopy, and X-ray diffraction.
Samples that were cut out of the Alloy 14 sheets were metallography polished in stages down to 0.02 μm grit to ensure smooth samples for scanning electron microscopy (SEM) analysis. SEM was done using a Zeiss EVO-MA10 model with the maximum operating voltage of 30 kV. Example SEM backscattered electron micrographs of the Alloy 14 sheet sample in the as-cast, HIPed and HIPed/heat treated conditions are shown in FIG. 30. The Alloy 14 sheet has a modal structure in as-cast state (FIG. 30a) where micron sized matrix grains are separated by lamella structure. The lamella structure can be clearly resolved in the as-cast sample by SEM. Alloy 14 as-cast sheet has a higher volume fraction of the lamella structure as compared to the Alloy 1 sheet (case Example #3) with larger lamella spacing. Additionally, evidence for austenite to ferrite transformation was found to occur during the casting in Alloy 14 sheet. The matrix grains are surrounded by a layer that appears to have different chemical composition according to the revealed contrast. The brighter edges of the grains indicate less B or Si content as compared to the darker grain interior resulting from the compositional element re-distribution during alloy solidification. After HIP cycle, the lamellas completely disappeared and were replaced by very fine precipitates distributed nearly homogeneous in the sample volume such that the matrix grain boundaries cannot be readily identified (FIG. 30b). After the heat treatment, some finer precipitates can be found in the sample (FIG. 30c).
Additional details of the Alloy 14 sheet structure are revealed using X-ray diffraction. X-ray diffraction was done using a Panalytical X′Pert MPD diffractometer with a Cu Kα x-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. The resulting scans were then subsequently analyzed using Rietveld analysis using Siroquant software. In FIGS. 31-33, X-ray diffraction scans are shown including the measured/experimental pattern and the Rietveld refined pattern for the Alloy 14 sheets in the as-cast, HIPed, and HIPed/heat treated conditions, respectively. As can be seen, good fits of the experimental data was obtained in all cases. Analysis of the X-ray patterns including specific phases found, their space groups and lattice parameters is shown in Table 13. Note that the space group represents a description of the symmetry of the crystal and can have one of 230 types and is further identified with its corresponding Hermann Maugin space group symbol.
In the as-cast sheet, three phases were identified, a cubic γ-Fe (austenite), a cubic α-Fe (ferrite) and a complex mixed transitional metal boride phase with the M2B stoichiometry. Note that the lattice parameters of the identified phases are different than that found for pure phases clearly indicating the dissolution of the alloying elements. For example, γ-Fe as a pure phase would exhibit a lattice parameter equal to a=3.575 Å, α-Fe would exhibit a lattice parameter equal to a=2.866 Å, and Fe2B1 pure phase would exhibit lattice parameters equal to a=5.099 Å and c=4.240 Å. Note that based on the significant change in lattice parameters in the M2B phase is it likely that silicon is also dissolved into this structure so it is not a pure boride phase. Additionally, as can be seen in Table 13, while the phases do not change, the lattice parameters do change as a function of the sheet condition (i.e. as-cast, HIPed, HIPed/heat treated), which indicates that redistribution of alloying elements is occurring.







TABLE 13







Rietveld Phase Analysis of Alloy 14 Sheet










Condition
Phase 1
Phase 2
Phase 3





As-Cast Sheet
γ-Fe
α-Fe
M2B



Structure: 
Structure: 
Structure:



Cubic
Cubic
Tetragonal



Space group #:
Space group #:
Space group #:



#225
#229
#140



Space group:
Space group:
Space group:



Fm3m
Im3m
14/mcm



LP: 
LP: 
LP: a = 5.156 Å



a = 3.589 Å
a = 2.880 Å
c = 4.240 Å


HIPed at 1000° C. 
γ-Fe
α-Fe
M2B


for 1 hour
Structure: 
Structure: 
Structure:



Cubic
Cubic
Tetragonal



Space group #:
Space group #:
Space group #:



#225
#229
#140



Space group:
Space group:
Space group:



Fm3m
Im3m
I4/mcm



LP: 
LP: 
LP: a = 5.275 Å



a = 3.587 Å
a = 2.862 Å
c = 4.003 Å


HIPed at 1000° C.
γ-Fe
α-Fe
M2B


for 1 hour, Heat
Structure: 
Structure: 
Structure:


treated at 350° C.
Cubic
Cubic
Tetragonal


for 20 minutes
Space group #:
Space group #:
Space group #:



#225
#229
#140



Space group:
Space group:
Space group:



Fm3m
Im3m
I4/mcm



LP: 
LP: 
LP: a = 5.226 Å



a = 3.591 Å
a = 2.872 Å
c = 4.025 Å









To examine the structural features of the Alloy 14 sheets in more details, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM samples, specimens were cut from the as-cast, HIPed, and HIPed/heat-treated sheets, and then ground and polished to a thickness of  ̃30 to  ̃40 μm. Discs were then punched from these polished thin sheets, and then finally thinned by twin-jet electropolishing for TEM observation. The microstructure examination was conducted in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV.
In FIG. 34, TEM micrographs of the microstructure of the Alloy 14 sheets in the as-cast, HIPed, and HIPed/heat treated sheets are shown. In the as-cast sample, the lamella structure is predominant (FIG. 34a) that is consistent with the SEM observation. The matrix grains are mostly less than 10 μm in size. Similar to SEM observations, the edge of the grains exhibits a different composition as compared to the grain interior. As shown in FIG. 34a, the TEM analysis also shows a layer around the matrix grain. This layer does not belong to the lamella structure as shown by the dash line. After HIP cycle, the lamella structure disappears, and is instead replaced with precipitates in the intergranular regions (FIG. 34b). In addition, precipitation also occurred inside the matrix grains such that no matrix grain boundaries can be clearly seen. This is a significant microstructural difference from Alloy 1 sheet, in which no precipitates form within the matrix grains during HIP cycle. After additional heat treatment, another significant change in the microstructure was observed. As shown in FIG. 34c, there is a marked grain refinement in the sample resulting from the heat treatment and grains of  ̃200 to  ̃300 nm in size were formed. As revealed by X-ray diffraction, the austenite to ferrite transformation is activated, which led to the grain refinement in accordance with Step #2 (Mechanism #1 Static Nanophase Refinement) towards development of the NanoModal Structure (Step #3).
Case Example #6
Tensile Properties and Structural Changes in Alloy 14
The tensile properties of the steel sheet produced in this application will be sensitive to the specific structure and specific processing conditions that the sheet experiences. In FIG. 35, the tensile properties of Alloy 14 sheet representing a Class 2 steel are shown in the as-cast, HIPed (1000° C. for 1 hour) and HIPed (1000° C. for 1 hour)/heat treated (350° C. for 20 minutes) conditions. As can be seen, the as-cast sheet shows much lower ductility than the HIPed and the HIPed/heat treated samples. This increase in ductility can be attributed to both the reduction of macrodefects in the HIPed sheets and microstructural changes occurring in the modal structures of the HIPed or HIPed/heat treated sheet as discussed earlier in Case Example #5. Additionally, during the application of a stress during tensile testing it will be shown the structural changes which are occurring.
For the Alloy 14 sheet HIPed at 1000° C. for 1 hour, additional structural details were obtained through using X-ray diffraction which was done on both the undeformed sheet samples and the gage sections of the deformed tensile specimens. X-ray diffraction was specifically done using a Panalytical X′Pert MPD diffractometer with a Cu Kα X-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. In FIG. 36, X-ray diffractions patterns are shown for the Alloy 14 sheet HIPed at 1000° C. for 1 hour in both the undeformed sheet condition and the gage section of the tensile tested specimen cut out from the sheet. As can be readily seen, there are significant structural changes occurring during deformation with new phases formation as indicated by new peaks in the X-ray pattern. Peak shifts indicate that redistribution of alloying elements is occurring between the phases present in both samples.
The X-ray pattern for the deformed Alloy 14 tensile tested specimen (HIPed (1000° C. for 1 hour) was subsequently analyzed using Rietveld analysis using Siroquant software. As shown in FIG. 37, a close agreement was found between the measured and calculated patterns. In Table 14, the phases identified in the Alloy 14 undeformed sheet and in a gage section of tensile specimens are compared. As can be seen, the M2B phase exists in the sheet before and after tensile testing although the lattice parameters changed indicating that the amount of solute elements dissolved in this phases changed. Additionally, the γ-Fe phase existing in the undeformed Alloy 14 sheet no longer exists in the gage section of tensile tested specimen indicating that a phase transformation took place. Rietveld analysis of the undeformed sheet and tensile tested specimen indicates that the volume fraction of α-Fe content exhibited only a slight increase measured from  ̃28% to  ̃29%. This would indicate that the γ-Fe phase transformed into multiple phases including possibly α-Fe and at least two new previously unknown phases. As shown in Table 14, after deformation, two new previously unknown hexagonal phases have been identified. One newly identified hexagonal phase is representative of a dihexagonal pyramidal class and has a hexagonal P63mc space group (#186) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 38a. The other hexagonal phase is representative of a ditrigonal dipyramidal class and has a hexagonal P6bar2C space group (#190) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 38b. It is theorized based on the small crystal unit cell size that this phase is likely a silicon based phase possibly a previously unknown Si-B phase. Note that in the FIG. 38, key lattice planes are identified corresponding to significant Bragg diffraction peaks.







TABLE 14







Rietveld Phase Analysis of Alloy 14 Sheet; 


Before and After Tensile Testing











Condition
Phase 1
Phase 2
Phase 3
Phase 4





Sheet -
γ-Fe
α-Fe
M2B



HIPed
Structure: 
Structure:
Structure:



at 1000° C.
Cubic
Cubic
Tetragonal



for 1 hour -
Space 
Space 
Space 



Prior to
group #:
group #:
group #:



tensile
#225
#229
#140



testing
Space group:
Space group:
Space group:




Fm3m
Im3m
I4/mcm




LP: 
LP: 
LP: 




a = 3.587 Å
a = 2.862 Å
a = 5.275 Å






c = 4.003 Å



Sheet -
α-Fe
M2B
Hexagonal
Hexagonal


HIPed
Structure: 
Structure:
Phase 1 
Phase 2 


at 1000° C.
Cubic
Tetragonal
(new)
(new)


for 1 hour - 
Space
Space
Structure:
Structure:


After 
group #:
group #:
Hexagonal
Hexagonal


tensile
#229
#140
Space 
Space 


testing
Space 
Space 
group #:
group #:



group:
group:
#186
#190



Im3m
I4/mcm
Space group:
Space group:



LP: 
LP: 
P63mc
P62barC



a = 2.870 Å
a = 5.150 Å
LP: 
LP: 




c = 4.195 Å 
a = 2.856 Å
a = 4.999 Å





c = 6.087 Å
c = 11.350 Å









To examine the structural changes of the Alloy 14 sheets induced by tensile deformation, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM samples, they were cut from the gage section of the tensile tested specimens and polished to a thickness of  ̃30 to  ̃40 μm. Discs were punched from these polished thin sheets, and then finally thinned by twin-jet electropolishing for TEM observation. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV.
In FIG. 39, the microstructure of the gage section of the Alloy 14 sheet in HIPed conditions before and after the tensile deformation is shown. In the sample before tension, the precipitates are distributed in the matrix. Additionally, fine grains are shown in the sample due to the grain refinement induced by the phase transformation during the HIP cycle corresponding to Step #2 (Static Nanophase Refinement). Thus, NanoModal Structure (Step #3) was developed in the material prior to deformation. After the yield stress is exceeded, further grain refinement is developed with the continued transformation of austenite phase induced by the tensile deformation. According to X-ray analysis, the austenite phase transforms into multiple phases simultaneously including two new unidentified phases. As a result, grains of  ̃200 to  ̃300 nm in size can be widely observed in the sample. Dislocation activity induced by tensile deformation can also be observed in some of the grains. At the same time, the boride precipitates retain the same geometry, suggesting that they do not experience obvious plastic deformation.
FIG. 40 shows a detailed microstructure of the gage section of the Alloy 14 sheet in HIPed conditions after the tensile deformation. In the microstructure, other than the hard boride phase exhibiting twinned structure, small grains of several hundred nanometers in size can be found. Moreover, the ring pattern of the electron diffraction pattern, which is a collective contribution from many grains, further confirms the refined microstructure. In the dark-field image, the small grains appear bright; their sizes are all less than 500 nm. Additionally, it can be seen that sub-structures are displayed within these small grains, indicating that the deformation-induced defects such as dislocations distort the lattice. As in Alloy 1, new hexagonal phases were identified in the sample after tensile deformation, which is believed to be the very fine precipitates that formed during the tensile deformation. Grain refinement might be considered as a result of Dynamic Nanophase Strengthening (Step #4) leading to High Strength NanoModal Structure (Step #5) in the Alloy 14 sheet.
As it was shown, the Alloy 14 sheet has demonstrated Structure #1 Modal Structure (Step#1) in as-cast state (FIG. 30a). High strength with high ductility in this material was measured after HIP cycle (FIG. 35), which provides the Static Nanophase Refinement (Step #2) and the formation of the NanoModal Structure (Step #3) in the material prior deformation. The strain hardening behavior of the Alloy 14 during tensile deformation is attributed mostly to grain refinement corresponding to Mechanism #2 Dynamic Nanophase Strengthening (Step #4) with subsequent creation of the High Strength NanoModal Structure (Step #5). Additional hardening may occur by dislocation mechanism in newly formed grains. The Alloy 14 sheet is an example of Class 2 steel with High Strength NanoModal Structure formation leading to high ductility at high strength.
Case Example #7
Structure Development in Alloy 19
According to the alloy stoichiometries in Table 2, the Alloy 19 was weighed out from high purity elemental charges. Similar to Alloy 14, this alloy has demonstrated Class 2 behavior with high plastic ductility at high strength. The resulting charges were arc-melted into 4 thirty-five gram ingots and flipped and remelted several times to ensure homogeneity. The resulting ingots were then re-melted and cast into 3 sheets under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. An example picture of one of the 1.8 mm thick Alloy 19 sheets is shown in FIG. 41. Two of the sheets were then HIPed at 1100° C. for 1 hour. One of the HIPed sheets was then subsequently heat treated at 700° C. for 20 minutes. The sheets in the as-cast, HIPed and HIPed/heat treated states were then cut up using a wire-EDM to produce samples for various studies including tensile testing, SEM microscopy, TEM microscopy, and X-ray diffraction.
Samples that were cut out of the Alloy 19 sheets were metallography polished in stages down to 0.02 μm grit to ensure smooth samples for scanning electron microscopy (SEM) analysis. The samples were analyzed in detail using a Zeiss EVO-MA10 model with the maximum operating voltage of 30 kV. Example SEM backscattered electron micrographs of the Alloy 19 sheet samples in the as-cast, HIPed and HIPed/heat treated conditions are shown in FIG. 42.
As shown in FIG. 42a, the microstructure of the as-cast Alloy 19 sheet distinctly exhibit modal structures, i.e., matrix grained phase and intergranular regions. The matrix grains are  ̃5 to  ̃10 μm in the size. Similar to the microstructure of Alloy 14, the edge of the grains exhibits different compositional contrast from that in the grain interior, perhaps due to the phase transformation during the casting. No lamella structure was revealed by SEM in as-cast state. Exposure to the HIP cycle led to significant changes in the microstructure. Very fine precipitates were formed that were nearly homogeneous distributed in the matrix grains and the intergranular regions so that the matrix grain boundaries cannot be readily identified (FIG. 42b). After the heat treatment, the volume fraction of precipitates increased significantly (FIG. 42c), most of which form with reduced microstructural scale.
Additional details of the Alloy 19 sheet structure are revealed using X-ray diffraction. X-ray diffraction was done using a Panalytical X′Pert MPD diffractometer with a Cu Kα X-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. The resulting scan patterns were then subsequently analyzed using Rietveld analysis using Siroquant software. In FIGS. 43-45, X-ray diffraction scan patterns are shown including the measured/experimental pattern and the Rietveld refined pattern for the Alloy 19 sheets in the as-cast, HIPed, and HIPed/heat treated conditions, respectively. As can be seen, good fits of the experimental data was obtained in all cases. Analysis of the X-ray patterns including specific phases found, their space groups and lattice parameters is shown in Table 15. Note that the space group represents a description of the symmetry of the crystal and can have one of 230 types and is further identified with its corresponding Hermann Maugin space group symbol.
In the as-cast sheet, three phases were identified, a cubic γ-Fe (austenite), a cubic α-Fe (ferrite) and a complex mixed transitional metal boride phase with the M2B stoichiometry. Note that the lattice parameters of the identified phases are different than that found for pure phases clearly indicating the dissolution of the alloying elements. For example, γ-Fe as a pure phase would exhibit a lattice parameter equal to a=3.575 Å, α-Fe would exhibit a lattice parameter equal to a=2.866 Å, and Fe2B1 pure phase would exhibit lattice parameters equal to a=5.099 Å and c=4.240 Å. Note that based on the significant change in lattice parameters in the M2B phase is it likely that silicon is also dissolved into this structure so it is not a pure boride phase. Additionally, as can be seen in Table 15, while the phases do not change, the lattice parameters do change as a function of the condition of the sheet (i.e. cast, HIPed, HIPed/heat treated) which indicates that redistribution of alloying elements is occurring.







TABLE 15







Rietveld Phase Analysis of Alloy 19 Sheet










Condition
Phase 1
Phase 2
Phase 3





As-Cast
γ-Fe
α-Fe
M2B



Structure: Cubic 
Structure: Cubic
Structure:



Space group #:
Space group #:
Tetragonal



#225
#229
Space group #:



Space group:
Space group:
#140



Fm3m
Im3m
Space group:



LP: a = 3.590 Å
LP: a = 2.868Å
I4/mcm





LP: a = 5.162 Å





c = 4.281 Å


HIPed at 1100° C.
γ-Fe
α-Fe
M2B


for 1 hour
Structure: Cubic 
Structure: Cubic
Structure:



Space group #:
Space group #:
Tetragonal



#225
#229
Space group #:



Space group:
Space group:
#140



Fm3m
Im3m
Space group:



LP: a = 3.593 Å
LP: a = 2.876 Å
I4/mcm





LP: a = 5.168 Å





c = 4.188 Å


HIPed at 1100° C.
γ-Fe
α-Fe
M2B


for 1 hour 
Structure: Cubic 
Structure: Cubic
Structure:


and heat treated 
Space group #:
Space group #:
Tetragonal


at 700° C.
#225
#229
Space group #:


for 20 minutes
Space group:
Space group:
#140



Fm3m
Im3m
Space group:



LP: a = 3.590 Å
LP: a = 2.873 Å
I4/mcm





LP: a = 5.197





c = 4.280









To examine the structural features of the Alloy 19 sheets in more details, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM samples, specimens were cut from the as-cast, HIPed, and HIPed/heat-treated sheets, and then ground and polished. To study the deformation mechanisms, samples were also taken from the gage section of the tensile tested specimens and polished to a thickness of  ̃30 to  ̃40 μm. Discs were punched from these polished thin sheets, and then finally thinned by twin-jet electropolishing for TEM observation. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope (TEM) operated at 200 kV.
In FIG. 46, TEM micrographs of the microstructure of the Alloy 19 sheets in the as-cast, HIPed, and HIPed/heat treated sheets are shown. In the as-cast sample, the grains of  ̃5 to  ̃10 μm in size with the lamella structure in the intergranular regions were observed (FIG. 46a). The lamella structure is much finer as compared to that in Alloy 14 sheets and was not previously revealed by SEM analysis. After the HIP cycle, the lamella structure generally disappears, and is instead replaced with precipitates that are homogeneously distributed in the sample volume (FIG. 46b). In addition, the refined grains can be observed after HIP cycle. The grain refinement is achieved through the phase transformation of austenite phase. As revealed by X-ray diffraction, the austenite to ferrite transformation is activated, which led to the grain refinement in accordance with Step #2 (Mechanism #1 Static Nanophase Refinement). After the heat treatment cycle, further grain refinement occurred as a result of the continued phase transformation resulting in the completion of the formation of the NanoModal Structure (Step #3). In addition, the precipitates become more uniformly distributed (FIG. 46c).
Case Example #8
Tensile Properties and Structural Changes in Alloy 19
The tensile properties of the steel sheet produced in this application will be sensitive to the specific structure and specific processing conditions that the sheet experiences. In FIG. 47, the tensile properties of Alloy 19 sheet representing a Class 2 steel are shown which were in the as-cast, HIPed (1100° C. for 1 hour), and HIPed (1100° C. for 1 hour)/heat treated (700° C. for 20 minutes) conditions. As can be seen, the as-cast sheet shows much lower ductility than the HIPed samples. This increase in ductility can be attributed to both the reduction of macrodefects in the HIPed sheets and microstructural changes occurring in the modal structures of the HIPed or HIPed/heat treated sheet as discussed earlier in Case Example #7. Additionally, during the application of a stress during tensile testing it will be shown that structural changes are occurring.
For the Alloy 19 sheet HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 20 minutes, additional structural details were obtained through using X-ray diffraction which was done on both the undeformed sheet samples and the gage sections of the deformed tensile specimens cut from the sheet. X-ray diffraction was specifically done using a Panalytical X′Pert MPD diffractometer with a Cu Kα x-ray tube and operated at 40 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. In FIG. 48, X-ray diffraction curves are shown of the Alloy 19 sheet HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 20 minutes for both the undeformed sheet and the gage section of tensile specimen from the same sheet after tensile deformation. As can be readily seen, there are significant structural changes occurring during deformation with new phases formation as indicated by new peaks in the X-ray pattern. Peak shifts indicate that redistribution of alloying elements is occurring between the phases present in both samples.
The X-ray pattern for the tensile tested specimen from Alloy 19 sheet (HIPed at 1100° C. for 1 hour and heat treated 700° C. for 20 minutes) was subsequently analyzed using Rietveld analysis using Siroquant software. As shown in FIG. 49, a close agreement was found between the measured and calculated patterns. In Table 16, the phases identified in the Alloy 19 undeformed sheet and a gage section of tensile specimens are compared. As can be seen, the M2B phase exists in the sheet before and after tensile testing although the lattice parameters changed indicating that the amount of solute elements dissolved changed. Additionally, the γ-Fe phase existing in the undeformed Alloy 19 sheet no longer exists in the tensile specimen gage section indicating that the phase transformation took place. Rietveld analysis of the undeformed sheet and tensile tested specimen indicates that the α-Fe content changes little with only a slight increase measured from  ̃65% to  ̃66%. This would indicate that the γ-Fe phase transformed into multiple phases including possibly α-Fe and at least two new previously unknown phases. As shown in Table 16, after deformation, two new previously unknown hexagonal phases have been identified. One newly identified hexagonal phase is representative of a dihexagonal pyramidal class and has a hexagonal P63mc space group (#186) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 50a. The other hexagonal phase is representative of a ditrigonal dipyramidal class and has a hexagonal P6bar2C space group (#190) and the calculated diffraction pattern with the diffracting planes listed is shown in FIG. 50b. It is theorized based on the small crystal unit cell size that this phase is likely a silicon based phase possibly a previously unknown Si-B phase. Note that in the FIG. 50, key lattice planes are identified corresponding to significant Bragg diffraction peaks.







TABLE 16







Rietveld Phase Analysis of Alloy 19 Sheet; 


Before and After Tensile Testing











Condition
Phase 1
Phase 2
Phase 3
Phase 4





Sheet - HIPed
γ-Fe
α-Fe
M2B



at 1000° C. for
Structure:
Structure:
Structure:



1 hour and
Cubic
Cubic
Tetragonal



heat treated at
Space group
Space group
Space group



700° C. for 20
#:
#:
#:



minutes - Prior
#225
#229
#140



to tensile
Space group:
Space group:
Space group:



testing
Fm3m
Im3m
I4/mcm




LP: 
LP: 
LP: 




a = 3.590 Å
a = 2.873 Å
a = 5.197






c = 4.280



Sheet - HIPed
α-Fe
M2B
Hexagonal
Hexagonal


at 1000° C. for
Structure:
Structure:
Phase 1 
Phase 2 


1 hour and
Cubic
Tetragonal
(new)
(new)


heat treated at
Space group
Space group
Structure:
Structure:


700° C. for 20
#:
#:
Hexagonal
Hexagonal


minutes - After
#229
#140
Space group 
Space group 


tensile testing
Space group:
Space group:
#:
#:



Im3m
I4/mcm
#186
#190



LP: 
LP: 
Space group:
Space group:



a = 2.865 Å
a = 5.086 Å
P63mc
P62barC




c = 4.206 Å
LP:
LP: 





a = 2.876 Å
a = 5.010 Å





c = 6.123 Å
c = 11.395 Å









To examine the structural changes of the Alloy 19 sheets induced by tensile deformation, high resolution transmission electron microscopy (TEM) was utilized to analyze the sample gage section before and after tensile tests. To prepare TEM sample, specimens were cut from the gage section of tensile specimens, and then ground and polished to a thickness of  ̃30 to  ̃40 μm. Discs were punched from these polished thin sheets, and then finally thinned by twin-jet electropolishing for TEM observation. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope (TEM) operated at 200 kV.
FIG. 51 shows TEM micrographs of microstructure in Alloy 19 sheet before and after the tensile deformation. As in Alloy 14, homogeneously distributed boride phase is found in the sample, and the austenite phase transformation during HIP cycle and heat treatment led to significant grain refinement as a result of Static Nanophase Refinement (Step #2) with NanoModal Structure (Step #3) in the sheet sample before deformation (FIG. 51a). In the sample after tensile testing, although the boride phase does not exhibit obvious plastic deformation, a significant structure change was observed that was induced by the deformation (FIG. 51b). First, many small grains of several hundred nanometers in size can be found. The electron diffraction in the inset of FIG. 51b shows the ring pattern, which shows the refinement in microstructure scale. The small grains can also be revealed in the dark-field image, as shown in FIG. 52, and the small grains less than 500 nm can be clearly seen. In addition, it can be found that the grains contain a high density of dislocations after the tensile deformation such that the lattice of many grains are distorted and appear as if they are further divided into smaller grains (FIG. 52b). FIG. 53 shows another example of TEM micrographs representing microstructure in the gage section of the tensile deformed sample. A number of dislocations generated in the grains can be seen, as indicated by the black arrows. In addition, nanometer size precipitates can be found in the microstructure, as indicated by the white arrows. These very fine precipitates are presumably the new phases induced by deformation and found in the X-ray diffraction scans. Fine grain formation is a result of Dynamic Nanophase Strengthening (Step #4) occurring in the sample during tensile deformation that leads to High Strength NanoModal Structure (Step #5) in the Alloy 19 sheet material.
As a summary, the deformation of Alloy 19 sheet is characterized by the substantial work hardening similar to that in Alloy 14 sheet. As it was shown, the Alloy 19 sheet has demonstrated Structure #1 Modal Structure (Step#1) in as-cast state (FIG. 46a). High strength with high ductility in this material was measured after HIP cycle and heat treatment, which provide the Static Nanophase Refinement (Step #2) and creation of the NanoModal Structure (Step #3) in the material prior deformation (FIG. 46c). The strain hardening behavior of the Alloy 19 during tensile deformation (FIG. 47) is attributed mostly to the previous grain refinement corresponding to Mechanism #2 Dynamic Nanophase Strengthening (Step #4) with subsequent High Strength NanoModal Structure (Step #5) represented in FIG. 51b and FIGS. 52-53. Additional hardening may occur by dislocation based mechanisms in newly formed grains. The Alloy 19 sheet is an example of Class 2 steel with High Strength NanoModal Structure formation leading to high ductility at high strength.
Case Example #9
Strain Hardening Behavior
Using high purity elements, 35 g alloy feedstocks of the targeted alloys listed in Table 2 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a PVC chamber, melted using RF induction and then ejected onto a copper die designed for casting a 3×4 inches plate with thickness of 1.8 mm. The resultant plates were subjected to HIP cycle with subsequent heat treatment. Corresponding HIP cycle parameters and heat treatment parameters are listed in Table 17. In a case of air cooling, the specimens were hold at the target temperature for a target period of time, removed from the furnace and cooled down in air. In a case of slow cooling, after the specimens were hold at the target temperature for a target period of time, the furnace was turned off and the specimens were cooled down with the furnace.
The listed samples from selected alloys (Table 17) were tested in tension on an Instron mechanical testing frame (Model 3369) with recording strain hardening coefficient values as a function of straining during testing utilizing Instron's Bluehill control and analysis software. The results are summarized in FIG. 54 where the strain hardening coefficient values are plotted versus corresponding plastic strain as a percentage of total elongation of the sample. As it can be seen, Samples 4 and 7 have demonstrated an increase in strain hardening after about 25% up to 80-90% of strain in the sample (FIG. 54a). These sheet samples have shown high ductility during tensile testing (FIG. 54b) and represents Class 1 steels. Sample 5 also represents Class 1 steels and demonstrated high ductility during tensile testing while strain hardening is almost independent from strain percentage with slight increase up to sample failure. For all these three samples, the strain hardening related to deformation of Modal Structure through dislocation mechanism with additional strengthening through Dynamic Nanophase Strengthening. Samples 1, 2 and 3 had demonstrated very high strain hardening at the strain value of about 50% with subsequent strain hardening coefficient values decreasing up to sample failure (FIG. 54a). These sheet samples have high strength/high ductility combination (FIG. 54b) and represents Class 2 steels where initial 50% of straining corresponds to phase transformation in the sample with a plateau on the stress-strain curve. Following strain hardening behavior corresponds to High Strength NanoModal Structure formation through extensive Dynamic Nanophase Strengthening. Sample 6 represents Class 2 steel also but have shown intermediate behavior in terms of strain hardening and intermediate properties at tensile testing that can be related to the lower level of phase transformation during straining depending on alloy chemistry.







TABLE 17







Sample Specification










Samples
Alloy
HIP Cycle
Heat Treatment





Sample 1
Alloy 24
1100° C. for 1 hour
None


Sample 2
Alloy 25
1100° C. for 1 hour
700° C. for 1 hour;





Slow cooling


Sample 3
Alloy 26
1100° C. for 1 hour
700° C. for 20 minutes;





Air cooling


Sample 4
Alloy 27
1100° C. for 1 hour
700° C. for 1 hour;





Air cooling


Sample 5
Alloy 28
1100° C. for 1 hour
700° C. for 1 hour;





Air cooling


Sample 6
Alloy 29
1100° C. for 1 hour
700° C. for 20 minutes;





Air cooling


Sample 7
Alloy 31
1100° C. for 1 hour
700° C. for 20 minutes;





Air cooling









Case Example #10
Strain Rate Sensitivity
Using high purity elements, 35 g alloy feedstocks of the Alloy 1 and Alloy 19 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a PVC chamber, melted using RF induction and then ejected onto a copper die designed for casting a 3×4 inches plate with thickness of 1.8 mm.
The resultant sheets from each alloy were subjected to HIP cycle using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time. The resultant plates were subjected to HIP cycle with subsequent heat treatment. Corresponding HIP cycle parameters and heat treatment parameters are listed in Table 18. In a case of air cooling, the specimens were hold at the target temperature for a target period of time, removed from the furnace and cooled down in air. In a case of slow cooling, after the specimens were hold at the target temperature for a target period of time, the furnace was turned off and the specimens were cooled down with the furnace.







TABLE 18







HIP Cycle and Heat Treatment Parameters











Alloy
HIP Cycle
Heat Treatment







Alloy 1
1000° C. for 1 hour
350° C. for 20 minutes;





Air cooling



Alloy 19
1125° C. for 1 hour
700° C. for 1 hour;





Slow cooling










The tensile measurements were done at four different strain rates on an Instron mechanical testing frame (Model 3369) utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving with the load cell attached to the top fixture. The displacement rate was varied in a range from 0.006 to 0.048 mm/sec. The resultant stress—strain curves are shown in FIGS. 55-56. Alloy 1 did not show strain rate sensitivity in a range of applied strain rates. Alloy 19 has demonstrated slightly higher strain hardening rate at lower strain rates in the studied range that is probably related to the volume fraction of dynamically refined phases induced by deformation at different strain rates.
Case Example #11
Sheet Material Behavior at Incremental Straining
Using high purity elements, 35 g alloy feedstocks of the Alloy 19 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a PVC chamber, melted using RF induction and then ejected onto a copper die designed for casting a 3×4 inches plate with thickness of 1.8 mm.
The resultant sheets from each alloy were subjected to HIP cycle at 1150° C. for 1 hour using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for 1 hour before cooling down to room temperature while in the machine.
The incremental tensile testing was done on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving while the load cell is attached to the top fixture. Each loading-unloading cycle was done at incremental strain of about 2%. The resultant stress—strain curves are shown in FIG. 57. As it can be seen, Alloy 19 has demonstrated strengthening at each loading-unloading cycle confirming Dynamic Nanophase Strengthening in the alloy during deformation at each cycle.
Case Example #12
Annealing Effect on Property Recovering
Using high purity elements, 35 g alloy feedstocks of the Alloy 19 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a PVC chamber, melted using RF induction and then ejected onto a copper die designed for casting a 3×4 inches plate with thickness of 1.8 mm.
The resultant sheet from the Alloy 19 was subjected to a HIP cycle using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature of 1100° C. was reached and were exposed to an isostatic pressure of 30 ksi for 1 hour. Subsequent heat treatment at 700° C. for 1 hour with slow cooling was applied to the sheet after the HIP cycle.
The tensile testing was done on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving with the load cell attached to the top fixture. Two tensile specimens were pre-strained to 10% with subsequent unloading. One of the samples was tested again up to failure. The resultant stress-strain curves are shown in FIG. 58a. As it can be seen, the Alloy 19 sheet after pre-straining has demonstrated high strength with limited ductility (−4.5%). Ultimate strength of the sample and summary strain from two tests correspond to that measured for the Alloy 19 sheets in the same conditions (same HIP cycle and heat treatment parameters) (see FIG. 57).
Another sample after pre-straining was annealed at 1150° C. for 1 hour with slow cooling and tested again up to failure. The resultant stress-strain curves are shown in FIG. 58b. The sample has demonstrated complete property restoration after annealing showing typical behavior of the Alloy 19 sheets in the same conditions (same HIP cycle and heat treatment parameters) without pre-straining (FIG. 47b).
Case Example #13
Cyclic Annealing Effect on Tensile Mechanisms
Using the methodology provided in Case Example #12 to prepare the sheet, an additional sample has been cut from Alloy 19 sheet after HIP cycle at 1100° C. for 1 hour and heat treatment at 700° C. for 1 hour. The sample was pre-strained to 10% with subsequent annealing at 1150° C. for 1 hour. Then it was deformed to 10% again with subsequent unloading and annealing at 1150° C. for 1 hour. This procedure was repeated 11 times total leading to total strain of  ̃100%. The tensile curves superimposed upon each other for all 11 cycles are shown in FIG. 59. The specimen after 10 cycles is shown in FIG. 60 as compared to its initial geometry. Note that same level of strength was recorded at each test cycle confirming property restoration at the annealing between tests.
High strength in pre-strained specimen (FIG. 58a) might be explained by High Strength Modal Structure Creation (Structure #3) during Dynamic Nanophase Strengthening (Mechanism #2) at tension. The restoration of the pre-strained sheet properties after annealing suggests that phase transformation at Dynamic Nanophase Strengthening (Mechanism #2) are reversible at subsequent annealing of the deformed material.
Microstructure of the gage section of the tensile specimens from Alloy 19 sheet (HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 1 hour) after pre-straining and after pre-straining with subsequent annealing was examined by scanning electron microscopy (SEM) using an EVO-60 scanning electron microscope manufactured by Carl Zeiss SMT Inc. Microstructure of the gage section of the tensile specimens from Alloy 19 sheet (HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 1 hour) after pre-straining to 10% is shown in FIG. 61. In the pre-strained microstructure (FIG. 61), no visible changes in microstructure have been revealed by SEM as compared to the Alloy 19 sheet before pre-straining (FIG. 42c). In a case of annealing at 1150° C. for 1 hour after pre-straining to 10%, the precipitates distribute even more homogeneously in the matrix (FIG. 62). Presumably some austenite is in the sample after annealing, but the austenite grains cannot be revealed. Due to the repetitive straining and annealing, this resulting microstructure may be considered as a prototype microstructure for future hot working like hot rolling.
Case Example #13
Bake Hardening of Sheet Material
Three by four inch plates with thickness of 1.8 mm were cast from Alloys 1, 2, and 3 with chemical composition specified in Table 2. The resultant sheets were subjected to HIP cycle using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature of 1100° C. was reached and were exposed to an isostatic pressure of 30 ksi for 1 hour. After the HIP cycle, the individual sheets were subsequently heat treated in a box furnace at 350° C. for 20 minutes. To evaluate the bake hardening effect, the resultant sheets were additionally annealed at 170° C. for 30 minutes.
Hardness measurements of sheet materials before and after bake hardening treatment were performed by Rockwell C Hardness test in accordance with ASTM E-18 standards. A Newage model AT130RDB instrument was used for all hardness testing which was done on  ̃9 mm by  ̃9 mm square samples cut from cast and treated sheets with thickness of 1.8 mm. Testing was done with indents spaced such that the distance between each of them was greater than three times the indent width. Hardness data (average of three measurements) for sheet materials before and after bake hardening treatment are listed in Table 19. As it can be seen, hardness increased in all three alloys after additional annealing demonstrating a favorable bake hardening effect in all three alloys.







TABLE 19







Bake Hardening Effect on Selected Alloys










HRC











(Average)











Alloy
Before
After
Bake Hardening Effect (Δ HRC)





Alloy 1
18.6
25.0
6.4


Alloy 2
23.8 
27.1
3.2


Alloy 3
21.9 
25.3
3.3









Case Example #15
Cold Formability of Sheet Material
A 3×4 inches plates with thickness of 1.8 mm were cast from Alloy 1, Alloy 2, and Alloy 3 with chemical composition specified in Table 2. The resultant sheets were subjected to HIP cycle using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time in accordance with Hc HIP cycle parameters listed in Table 6. Resultant sheets were subjected to Erichsen Cup Test (ASTM E643-09) to estimate cold formability of the cast sheet materials. The Erichsen cupping test is a simple stretch forming test of a sheet clamped firmly between blank holders to prevent in-flow of sheet material into the deformation zone. The punch is forced onto the clamped sheet with tool contact (lubricated, but with some friction) until cracks occur. The depth (mm) of the punch is measured and gives the Erichsen depth index as shown in FIG. 63. Test results for sheets from selected alloys are listed in Table 20 showing variation in depth index from 2.72 to 5.48 mm depending on alloy chemistry. These measurements correspond to plastic ductility of the plate at outer surface in a range from 9 to 20% indicating significant plasticity of the selected alloys.







TABLE 20







Erichsen Cup Test Results for As-Cast Plates












Maximum 
Erichsen




Load
depth index



Alloy 
(kN)
(mm)







Alloy 1
9.00
5.18



Alloy 2
9.72
2.72



Alloy 3
8.15
5.48










The selected three alloys represent deformation behavior corresponding to that described in Case Example #4 when only Step #1 (Modal Structure) and Step #4 (Dynamic Nanophase Strengthening) was observed. High levels of formability might be achieved in the alloys with referenced chemistries that demonstrate deformation behavior described in Case Examples #6 and #8. Due to Static Nanophase Refinement (Step #2) and NanoModal Structure (Step #3), a reversible phase transformation with Dynamic Nanophase Strengthening (Step #4) was found as described in Case Example #12. By applying annealing to pre-deformed sheet material, total strain of more than 100% might be achieved.
Case Example #16
Thick Plate Properties
Using high purity elements, feedstocks with different mass of the Alloy 1 and Alloy 19 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the crucible of a custom-made vacuum casting system. The feedstock was melted using RF induction and then ejected onto a copper die designed for casting a 4×5 inches sheets at different thickness. Sheets with three different thicknesses of 0.5 inches, 1 inch and 1.25 inches were cast from each alloy (FIG. 64). Note that the sheets that were cast were much thicker than the previous 1.8 mm plates and illustrate the potential for the chemistries in Table 2 to be processed by the Thin Slab Casting process.
All sheets from each alloy were subjected to HIP cycle using an American Isostatic Press Model 645 machine with a molybdenum furnace with furnace chamber size of 4 inch diameter by 5 inch height. The sheets were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time. HIP cycle parameters for both alloys are listed in Table 21 and are representative of the thermal exposure experienced by sheets in the Thin Slab Casting process. After HIP cycle, sheet material was heat treated in a box furnace at parameters specified in Table 22







TABLE 21







HIP Cycle Parameters













HIP Cycle
HIP Cycle
HIP Cycle




Temperature
Pressure
Time



Alloy
[° C.]
[psi]
[hr]







Alloy 1
1000
30,000
1



Alloy 19
1125
30,000
1

















TABLE 22







Heat Treatment Parameters













Temperature
Time




Alloy
(° C.) 
(min)
Cooling







Alloy 1
350
20
In air



Alloy 19
700
60
With furnace










The tensile specimens were cut from the sheets using wire electrical discharge machining (EDM). The tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving with the load cell attached to the top fixture. In Table 23, a summary of the tensile test results including total tensile strain, yield stress, ultimate tensile strength and Elastic Modulus is shown for 1.25 inches thick sheets in as-cast state and after HIP cycle and heat treatment. As can be seen the tensile strength values vary from 428 to 575 MPa for Alloy 1 sheet and from 642 to 814 MPa for Alloy 19 sheet. The total strain value varies from 2.78 to 14.20% for Alloy 1 sheet and from 3.16 to 6.02% for Alloy 19 sheet. Elastic Modulus is measured in a range from 103 to 188 GPa for both alloys. Note that these properties are not optimized at the much greater cast thickness but represent clear indications of the promise of the new steel types, enabling structures and mechanisms for large scale production through Thin Slab Casting.







TABLE 23







Summary of Tensile Test Results for 1.25 inches Thick Sheets













Sheet
Yield
Ultimate
Tensile
Elastic



Thickness
Stress
Strength
Elongation
Modulus


Alloy
(inches)
(MPa)
(MPa)
(%)
(GPa)















Alloy 1
As-cast
237
518
8.78
165




226
428
2.78
152




256
525
10.10
172




242
515
7.39
169




229
555
13.49
152




242
543
11.58
103



HIPed
234
575
14.20
165



and heat
222
496
6.78
124



treated
237
533
11.80
117


Alloy 19
As-cast
377
760
5.35
167




334
751
5.47
134




387
665
4.59
176




329
642
4.26
188




371
687
4.83
155




353
652
4.98
162



HIPed
318
805
6.02
150



and heat
344
814
5.96
153



treated
366
809
5.61
154




284
656
3.16
134









Case Example #17
Melt-Spinning Study
Using high purity elements, 15 g alloy feedstocks of the Alloy 19 were weighed out according to the atomic ratios provided in Table 2. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a melt-spinning chamber in a quartz crucible with a hole diameter of  ̃0.81 mm. The ingots were then processed by melting in different atmosphere using RF induction and then ejected onto a 245 mm diameter copper wheel which was traveling at different tangential velocities varying from 16 to 39 m/s. Continuous ribbons with various thicknesses were produced.
Thermal analysis was done on the as-solidified ribbon structure on a Perkin Elmer DTA-7 system with the DSC-7 option. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) was performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultrahigh purity argon. All ribbons have crystalline structure in as-cast state and similar melting behavior with melting peak at 1248° C.
The mechanical properties of metallic ribbons were obtained at room temperature using microscale tensile testing. The testing was carried out in a commercial tensile stage made by Fullam which was monitored and controlled by a MTEST Windows software program. The deformation was applied by a stepping motor through the gripping system while the load was measured by a load cell that was connected to the end of one gripping jaw. Displacement was obtained using a Linear Variable Differential Transformer (LVDT) which was attached to the two gripping jaws to measure the change of gauge length. Before testing, the thickness and width of a ribbon were carefully measured for at least three times at different locations in the gauge length. The average values were then recorded as gauge thickness and width, and used as input parameters for subsequent stress and strain calculation. The initial gauge length for tensile testing was set at  ̃9 mm with the exact value determined after the ribbon was fixed by accurately measuring the ribbon span between the front faces of the two gripping jaws. All tests were performed under displacement control, with a strain rate of  ̃0.001s−1. A summary of the tensile test results including total elongation, yield strength, ultimate tensile strength, and Young's Modulus are shown in Table 24. As can be seen the tensile strength values vary from 810 MPa to 1288 MPa with the total elongation from 0.83% to 17.33%. Large scattering in properties is observed for all tested ribbons suggesting a formation of non-uniform structures at fast cooling.







TABLE 24







Summary on Tensile Properties of Melt-Spun Ribbons










Wheel Speed
Yield Stress
Ultimate Strength
Total


(m/s)
(MPa)
(MPa)
Elongation (%)













16
664
829
9.82



665
810
2.17



701
828
5.61


20
799
891
3.72



769
922
9.89



733
1095
17.33


25
751
1020
15.56



1003
1142
2.51



746
1043
15.06


30
1113
1249
2.82



770
1027
15.67



1183
1288
1.39


39
1075
1220
1.13



650
837
0.83



1030
1193
1.14









Case Example #18
Tensile Properties of Mn-Containing Alloys
Tensile Properties of alloys listed in Table 25 were examined to determine the effect of the addition of Manganese in levels of up to 4.53 atomic percent. Alloys were prepared in 35 g charges using high purity research grade elemental constituents. Charges of each alloy were arc-melted into ingots, and then homogenized under argon atmosphere. The resulting 35 gram ingots were then cast into plates with nominal dimensions of 65 mm by 75 mm by 1.8 mm.







TABLE 25







Alloy Composition















Alloy
Fe
Cr
Ni
B
Si
Mn







Alloy 25
62.20
17.62
4.14
5.30 
6.60
4.14



Alloy 26
60.35
20.70
3.53
5.30
6.60
3.52



Alloy 27
61.10
19.21
3.90
5.30
6.60
3.89



Alloy 28
61.32
20.13
3.33
5.30
6.60
3.32



Alloy 29
63.83
17.97
3.15
5.30
6.60
3.15



Alloy 30
63.08
15.95
4.54
5.30
6.60
4.53



Alloy 31
64.93
16.92
3.13
5.30
6.60
3.12



Alloy 32
64.45
15.86
3.90
5.30 
6.60
3.89



Alloy 33
62.11
20.31
2.84
5.30
6.60
2.84



Alloy 34
62.20
17.62
6.21
5.30
6.60
2.07



Alloy 35
60.35
20.70
5.29
5.30
6.60
1.76



Alloy 36
61.10
19.21
5.85
5.30
6.60
1.94



Alloy 37
61.32
20.13
4.99
5.30
6.60
1.66



Alloy 38
63.83
17.97
4.73
5.30
6.60
1.57



Alloy 39
63.08
15.95
6.80
5.30
6.60
2.27



Alloy 40
64.93
16.92
4.69
5.30
6.60
1.56



Alloy 41
64.45
15.86
5.85
5.30
6.60
1.94



Alloy 42
62.11
20.31
4.26
5.30
6.60
1.42










As-cast plates were then subjected to hot isostatic pressing (HIPing) at 30 ksi for 1 hour, with a temperature selected according to Table 26. HIPing was done using an American Isostatic Press Model 645 machine with a molybdenum furnace. Samples were heated to the target temperature at a rate of 10° C./min and held at temperature under the pressure of 30 ksi for 1 hour.







TABLE 26







HIP Parameters Selected for Alloys Used in Case Study












HIP Cycle
HIP
HIP
Dwell


Alloy
Designation
Temperature
Pressure
Time





Alloy 25
Hf
1150° C.
30 ksi
1 Hour


Alloy 26
Hf
1150° C.
30 ksi
1 Hour


Alloy 27
Hf
1150° C.
30 ksi
1 Hour


Alloy 28
Hf
1150° C.
30 ksi
1 Hour


Alloy 29
Hf
1150° C.
30 ksi
1 Hour


Alloy 30
Hf
1150° C.
30 ksi
1 Hour


Alloy 31
Hf
1150° C.
30 ksi
1 Hour


Alloy 32
Hf
1150° C.
30 ksi
1 Hour


Alloy 33
Hf
1150° C.
30 ksi
1 Hour


Alloy 34
Hf
1150° C.
30 ksi
1 Hour


Alloy 35
Hf
1150° C.
30 ksi
1 Hour


Alloy 36
Hf
1150° C.
30 ksi
1 Hour


Alloy 37
Hf
1150° C.
30 ksi
1 Hour


Alloy 38
Hf
1150° C.
30 ksi
1 Hour


Alloy 39
Hf
1150° C.
30 ksi
1 Hour


Alloy 40
Hf
1150° C.
30 ksi
1 Hour


Alloy 41
Hf
1150° C.
30 ksi
1 Hour


Alloy 42
Hf
1150° C.
30 ksi
1 Hour









Tensile specimens were cut from HIPed plates by Electric Discharge Machining (EDM). Some of the tensile specimens were heat treated according to the heat treatment schedule in Table 27. Heat treatments were performed using a Lindberg Blue furnace. In a case of air cooling, the specimens were held at the target temperature for a target period of time, removed from the furnace and cooled down in air. In a case of slow cooling, the specimens were heated to the target temperature and then cooled with the furnace at cooling rate of 1° C./min. Heat treated specimens were then tested to determine tensile properties of the selected alloys.







TABLE 27







Heat Treatment Schedule for Case Study Alloys










Heat

Dwell



Treatment
Temperature
Time
Cooling





HT2
700° C.
1 Hour
Air Cooling


HT3
700° C.
N/A
1° C./min Slow Cool


HT4
850° C.
1 Hour
Air Cooling









Tensile testing was performed on an Instron Model 3369 mechanical testing frame, using the Instron Bluehill control and analysis software. Samples were tested at room temperature under displacement control at a strain rate of 1×10−3 per second. Samples were mounted to a stationary bottom fixture, and a top fixture attached to a moving crosshead. A 50 kN load cell was attached to the top fixture to measure load. Strain measurements were made using an advanced video extensometer (AVE). Tensile results for the study are tabulated in Table 28. As can be seen from the results table, tensile strength in the examined alloys ranged from 753 to 1511 MPa. It is useful to note that the ceramics used in the production of sheets for the indicated case examples (e.g. ceramic crucibles) were not optimized for these manganese containing melts. This resulted in some ceramic entrainment in the melt creating defects which lowered the ductility in some cases. Higher ductility is expected by changing the ceramics used in melting. Total elongation values ranged from 2.0% to 28.0%. Strain hardening exponents were calculated as an average value, using a strain range beginning with the yield point and ending with the point corresponding to the ultimate tensile strength. Example tensile curves have been provided in FIG. 65 showing variation in alloy mechanical response depending on alloy chemistry and processing conditions.







TABLE 28







Tensile Properties of Manganese Containing Alloys


















Yield
Ultimate
Tensile
Elastic
Strain
Type 



HIP
Heat
Stress 
Strength
Elongation
Modulus
Hardening
of


Alloy
Cycle
Treatment
(MPa)
(MPa)
(%)
(GPa)
Exponent
Behavior


















Alloy 25
Hf
None
472
1020
10.8
169
0.57
Class 2





473
914
9.8
213
0.54
Class 2





484
1045
11.5
183
0.56
Class 2




HT2
507
1244
14.4
183
0.69
Class 2





505
1247
13.9
184
0.71
Class 2




HT3
492
1204
13.2
177
0.70
Class 2





500
1076
10.7
187
0.65
Class 2




HT4
505
1095
12.2
150
0.62
Class 2





525
1288
16.8
174
0.69
Class 2


Alloy 26
Hf
None
651
1018
8.7
132
0.28
Class 2





642
990
7.4
187
0.25
Class 2




HT2
502
973
7.7
143
0.26
Class 2





624
846
4.6
192
0.14
Class 2




HT3
617
753
2.0
172
0.15
Class 1





616
889
4.8
279
0.13
Class 1




HT4
634
1151
14.9
200
0.32
Class 2


Alloy 27
Hf
None
585
1196
14.1
189
0.46
Class 2




HT2
548
1124
11.9
172
0.47
Class 2





567
1235
15.3
167
0.49
Class 2




HT3
582
1131
11.2
190
0.46
Class 2





611
983
8.1
175
0.32
Class 2




HT4
626
1200
18.2
161
0.41
Class 2





556
1098
11.4
177
0.41
Class 2


Alloy 28
Hf
None
552
779
2.7
223
0.20
Class 1





657
878
3.5
222
0.14
Class 1




HT2
648
1083
10.4
180
0.29
Class 2




HT3
671
846
2.1
207
0.16
Class 1





633
851
2.7
225
0.14
Class 1




HT4
601
1094
12.7
232
0.31
Class 2


Alloy 29
Hf
None
1038
1239
2.4
139
0.19
Class 2





573
996
2.5
184
0.38
Class 2




HT2
558
1254
10.7
162
0.37
Class 2




HT3
665
964
3.1
206
0.24
Class 2





702
1280
9.2
183
0.33
Class 2




HT4
556
1227
6.8
187
0.61
Class 2





573
1129
5.5
148
0.61
Class 2


Alloy 30
Hf
None
459
1203
13.0
155
0.82
Class 2





474
1341
17.7
126
0.82
Class 2





466
1275
14.3
153
0.82
Class 2




HT2
432
1348
18.3
148
0.80
Class 2





450
1323
16.2
160
0.85
Class 2




HT3
445
768
7.5
186
0.40
Class 2





448
1356
20.6
153
0.77
Class 2





425
1156
13.4
147
0.77
Class 2




HT4
437
1115
12.0
149
0.80
Class 2





420
1355
17.5
185
0.83
Class 2





429
1021
10.5
160
0.69
Class 2


Alloy 31 
Hf
None
650
1330
4.5
194
0.37
Class 2





676
1373
7.4
179
0.32
Class 2




HT2
661
1169
5.7
198
0.31
Class 2




HT3
732
973
2.7
204
0.18
Class 1





790
1011
2.7
239
0.15
Class 1




HT4
481
1160
4.0
184
0.47
Class 2





469
1139
4.6
174
0.55
Class 2





502
1245
6.0
163
0.49
Class 2


Alloy 32
Hf
None
432
1391
10.6
127
0.94
Class 2





454
1381
8.8
198
0.89
Class 2




HT2
431
1423
13.3
196
0.91
Class 2





418
1434
12.6
142
0.92
Class 2





366
872
5.4
160
0.67
Class 2




HT3
410
1390
9.6
153
0.94
Class 2





384
1421
13.2
149
0.90
Class 2





398
1418
9.4
152
0.95
Class 2




HT4
398
1444
15.8
155
0.92
Class 2





451
1431
13.9
187
0.97
Class 2





444
1349
9.9
155
0.98
Class 2


Alloy 33
Hf
None
657
1100
5.1
211
0.27
Class 1





743
1064
4.3
225
0.19
Class 1




HT2
701
1100
11.5
235
0.21
Class 1




HT3
749
1013
3.4
224
0.19
Class 1





680
983
2.8
243
0.22
Class 1




HT4
697
1080
7.6
238
0.20
Class 1


Alloy 34
Hf
None
440
1228
18.8
137
0.77
Class 2





438
1236
18.9
185
0.70
Class 2





449
1273
21.1
152
0.73
Class 2




HT2
418
1124
15.0
169
0.73
Class 2





438
1222
18.2
183
0.72
Class 2





430
1278
25.6
137
0.76
Class 2




HT3
435
1193
16.9
172
0.72
Class 2





421
1261
26.7
147
0.75
Class 2





426
1262
20.4
141
0.73
Class 2





460
1208
17.7
129
0.76
Class 2




HT4
425
1180
17.2
141
0.76
Class 2





426
1194
17.6
159
0.74
Class 2





443
1148
16.3
135
0.70
Class 2





460
1292
28.0
103
0.74
Class 2


Alloy 35
Hf
None
526
927
11.2
183
0.31
Class 2





580
1114
17.2
227
0.44
Class 2





583
1162
19.3
168
0.44
Class 2




HT2
501
1024
13.2
197
0.53
Class 2





518
978
12.1
186
0.48
Class 2





541
972
11.9
116
0.41
Class 2




HT3
564
856
8.0
185
0.26
Class 2





594
1095
14.6
195
0.45
Class 2





561
1047
12.8
219
0.43
Class 2




HT4
571
1168
18.4
194
0.49
Class 2





594
1046
12.6
176
0.43
Class 2





584
990
11.7
202
0.39
Class 2


Alloy 36
Hf
None
440
1210
20.8
155
0.70
Class 2





461
1169
18.2
193
0.68
Class 2




HT2
441
952
12.3
199
0.57
Class 2





435
1084
15.2
194
0.63
Class 2





472
1200
20.1
114
0.71
Class 2




HT3
412
996
13.5
258
0.60
Class 2





434
1205
23.1
176
0.68
Class 2





463
1029
14.3
149
0.60
Class 2




HT4
463
1243
27.1
126
0.67
Class 2





455
1166
18.9
131
0.69
Class 2





424
1194
19.7
192
0.71
Class 2





437
1243
26.7
194
0.66
Class 2


Alloy 37
Hf
None
539
1181
15.4
166
0.61
Class 2





563
1178
15.7
145
0.58
Class 2




HT2
541
1186
16.4
194
0.56
Class 2





510
1180
17.0
187
0.56
Class 2




HT3
542
1204
18.1
186
0.55
Class 2





503
1185
15.0
228
0.59
Class 2





519
1015
9.5
220
0.53
Class 2




HT4
523
1114
12.5
156
0.59
Class 2





582
1200
19.0
116
0.53
Class 2





553
1187
17.5
168
0.51
Class 2


Alloy 38
Hf
None
465
1319
9.0
157
0.84
Class 2





437
1275
8.1
256
0.88
Class 2




HT2
418
1347
12.9
127
0.82
Class 2





407
1304
11.3
182
0.94
Class 2




HT3
435
1279
6.1
157
0.85
Class 2





419
1289
13.3
184
0.80
Class 2





431
1312
11.9
185
0.81
Class 2




HT4
433
1354
10.6
139
0.99
Class 2





434
1342
12.5
181
0.95
Class 2


Alloy 39
Hf
None
454
787
8.9
204
0.44
Class 2





443
1065
14.3
166
0.68
Class 2





458
1132
16.1
177
0.70
Class 2




HT2
452
1011
12.6
190
0.66
Class 2





445
996
12.3
190
0.65
Class 2




HT3
457
1273
23.9
157
0.72
Class 2





448
1296
23.8
161
0.70
Class 2





446
1277
20.9
159
0.74
Class 2





424
1159
16.6
181
0.80
Class 2




HT4
466
1092
14.8
184
0.68
Class 2





437
1163
17.0
163
0.74
Class 2





444
954
12.1
180
0.60
Class 2


Alloy 40
Hf
None
661
1492
5.3
155
0.42
Class 2





669
1511
9.9
203
0.36
Class 2





673
1510
8.1
225
0.35
Class 2




HT2
617
1306
7.5
224
0.48
Class 2





638
1343
11.5
193
0.42
Class 2





648
1325
8.8
191
0.44
Class 2




HT3
802
1383
7.9
193
0.33
Class 2





830
1368
8.2
186
0.31
Class 2





830
1408
11.4
186
0.30
Class 2





815
1391
8.9
201
0.32
Class 2




HT4
416
1357
10.1
183
0.89
Class 2





402
1390
11.4
153
0.87
Class 2





401
1356
7.3
204
0.98
Class 2





425
1399
13.4
213
0.88
Class 2


Alloy 41
Hf
None
447
1372
13.7
161
0.49
Class 2





458
1029
8.9
155
0.37
Class 2




HT2
409
1150
8.7
164
0.95
Class 2





401
1372
16.4
150
0.88
Class 2




HT3
387
937
7.2
142
0.69
Class 2





395
1386
14.6
179
0.86
Class 2





394
1180
9.1
192
0.97
Class 2




HT4
441
1319
11.4
131
0.96
Class 2





446
810
6.9
132
0.74
Class 2





438
1366
14.9
123
0.98
Class 2


Alloy 42
Hf
None
583
1244
10.7
174
0.59
Class 2





596
924
5.6
164
0.38
Class 2




HT2
579
1188
8.1
179
0.56
Class 2





572
1202
9.8
213
0.54
Class 2





531
1135
7.0
246
0.61
Class 2




HT3
382
1171
7.8
172
0.66
Class 2





585
992
5.2
192
0.51
Class 2





625
1047
6.0
119
0.51
Class 2




HT4
593
1085
7.9
206
0.43
Class 2





574
1196
13.0
199
0.45
Class 2





619
779
3.5
193
0.17
Class 2









Case Example #19
Melt-Spinning Study on Additional Alloys
Melt-spinning is an example of chill surface processing in which high cooling rates, higher than either thin slab or twin-roll casting, may be achieved. The required charge size is small and the process is faster compared to the other formerly noted processes. Thus, it is useful tool for quickly examining the potential of an alloy for chill surface processing. Using high purity elements, 15 g charges of the alloys listed in Table 29 were weighed. Charges were then placed into the copper hearth of an arc-melting system. The charge was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a melt-spinning chamber in a quartz crucible with an orifice diameter of  ̃0.81 mm.







TABLE 29







Alloy Chemistries














Alloy
Fe
Cr
Ni
B
Si
Mn
C

















Alloy 43
62.38 
17.40 
7.92 
7.40
4.20
0.50
0.20


Alloy 44
65.99 
13.58 
6.58 
7.60
4.40
1.50
0.35


Alloy 45
58.76 
17.22 
9.77 
7.80
4.60
1.30
0.55


Alloy 46
58.95 
11.35
13.40
8.00
4.80
1.25
2.25


Alloy 47
62.28 
10.00
12.56
4.80
8.00
2.00
0.36


Alloy 48
53.82 
20.22
11.60
4.60
7.80
0.75
1.21


Alloy 49
61.21 
21.00
4.90 
4.40
7.60
0.00 
0.89


Alloy 50
62.00 
17.50
6.25 
4.20
7.40
0.10 
2.55


Alloy 51
59.71 
14.30
13.74 
4.00
7.20
0.40 
0.65


Alloy 52
57.85
13.90
12.25
7.00
7.00
1.75 
0.25


Alloy 53
56.90
15.25
14.50
6.00
6.00
1.35 
0.00


Alloy 54
65.82
12.22 
7.22 
5.00
6.00
1.14 
2.60


Alloy 55
58.72
18.26 
8.99 
4.26
7.22
1.55 
1.00


Alloy 56
61.30
17.30 
6.50 
7.15
4.55
0.20 
3.00


Alloy 57
65.80
14.89 
8.66 
4.35
4.05
0.00 
2.25


Alloy 58
63.99
12.89
10.25
8.00
4.22
0.65
0.00


Alloy 59
71.24
10.55
5.22
7.55
4.55
0.89
0.00


Alloy 60
61.88
11.22
12.55 
7.45
5.22
1.12 
0.56









The density of the alloys was measured on arc-melt ingots using the Archimedes method in a balance allowing weighing in both air and distilled water. The density of each alloy is tabulated in Table 30 and was found to vary from 7.45 g/cm3 to 7.71 g/cm3. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3.







TABLE 30







Summary of Density Results (g/cm3)










Alloy
Density (avg)







Alloy 43
7.66



Alloy 44
7.65



Alloy 45
7.63



Alloy 46
7.67



Alloy 47
7.62



Alloy 48
7.54



Alloy 49
7.45



Alloy 50
7.54



Alloy 51
7.64



Alloy 52
7.60



Alloy 53
7.67



Alloy 54
7.61



Alloy 55
7.57



Alloy 56
7.59



Alloy 57
7.66



Alloy 58
7.71



Alloy 59
7.54



Alloy 60
7.67










The arc-melted fingers were then placed into a melt-spinning chamber in a quartz crucible with a orifice diameter of  ̃0.81 mm. The ingots were then processed by melting in different atmosphere using RF induction and then ejected onto a 245 mm diameter copper wheel which was traveling at a tangential velocity at 20 m/s. Continuous ribbons with thicknesses between 41 μm and 59 μm were produced. The quality of ribbon produced varied by alloy with some alloys providing more uniform cross-sections than others.
Differential Thermal Analysis (DTA) was performed on the as-solidified ribbon using a Netzsch DSC 404 F3 Pegasus system. Scans were performed at a constant heating rate of 10° C./minute from 100° C. to 1410° C. with an ultrahigh purity argon purge gas to protect samples from oxidation as shown in Table 31. As shown, some ribbons (melt-spun at 20 m/s) contained small fractions of metallic glass while others did not. Based on the thickness of the ribbon produced, the estimated cooling rates were 3×105 to 6×105K/s which is beyond the cooling rates identified for sheet as described previously. For the alloys in this case example, melting was found to occur with one to three distinct melting peaks. The solidus ranged between 1138° C. and 1230° C. with melting events observed up to 1374° C.







TABLE 31







Differential Thermal Analysis Data for Melting Behavior













Metallic 







Glass
Solidus 
Peak 1 
Peak 2 
Peak 3


Alloy
Present
(° C.)
(° C.)
(° C.)
(° C.)















Alloy 43
No
1241
1256
1264
1271


Alloy 44
Yes
1221
1244
1250
—


Alloy 45
Yes
1227
1245
1260
1270


Alloy 46
Yes
1138
1155
1205
1218


Alloy 47
No
1185
1215
1241
1313


Alloy 48
No
1216
1252
—
—


Alloy 49
No
1208
1223
1273
—


Alloy 50
No
1180
1197
1218
—


Alloy 51
No
1218
1244
1302
1349


Alloy 52
Yes
1198
1215
1240
1245


Alloy 53
No
1221
1242
1248
1252


Alloy 54
No
1157
1173
—
—


Alloy 55
No
1230
1255
—
—


Alloy 56
Yes
1180
1198
1248
—


Alloy 57
No
1226
1250
1374
—


Alloy 58
Yes
1215
1238
1243
1251


Alloy 59
No
1211
1226
1240
—


Alloy 60
Yes
1193
1228
1236
1292









The mechanical properties of metallic ribbons were measured at room temperature using uniaxial tensile testing. The testing was carried out in a commercial tensile stage made by Fullam which was monitored and controlled by a MTEST Windows software program. Deformation was applied by a stepping motor through the gripping system while the load was measured by a load cell which was connected to the end of one gripping jaw. Displacement was measured using a Linear Variable Differential Transformer (LVDT) which was attached to the two gripping jaws to measure the change of gauge length. Before testing, the thickness and width of a ribbon were carefully measured for at least three times at different locations in the gauge length. The average values were then recorded as gauge thickness and width, and used as input parameters for subsequent stress and strain calculations. The initial gauge length for tensile testing was set at  ̃9 mm with the exact value determined after the ribbon was fixed by measuring the ribbon span between the front faces of the two gripping jaws.
All tests were performed under displacement control, with a strain rate of  ̃0.001s−1. Three tests were performed for each bendable ribbon while one to three tests were performed on non-bendable ribbons. A summary of the tensile test results including total elongation, yield strength, and ultimate tensile strength are shown in Table 32. The tensile strength values varied from 282 to 2072 MPa. The total elongation value varied from 0.37 to 6.56% indicating limited ductility of alloys in as-cast state for most samples. Some samples failure occurred in elastic region without yielding while others showed clear ductility such Alloy 47 shown in FIG. 66. Considerable variability exists in the mechanical properties of these ribbons as this variability is caused in part by irregularities in sample geometry and microstructural defects which means that the tensile properties are lower than expected in sheet form. Additionally, for alloys which contained metallic glass (i.e. 44, 45, 46, 52, 56, 58, and 60), it can be seen that the mechanical properties especially the ductility were lowered. Thus, it is clear that the favorable structures and mechanisms in this application are for crystalline structures and not partial or full metallic glass.







TABLE 32







Summary on Tensile Properties of 


Melt-Spun Ribbons at 20 m/s











Yield Stress
Ultimate Strength
Tensile Elongation


Alloy
(MPa)
(MPa)
(%)













Alloy 43
1663
2072
3.63



1225
1611
3.37



1241
1618
3.13


Alloy 44
904
1085
1.08


Alloy 45
282
282
0.37


Alloy 46
1958
2019
2.59


Alloy 47
630
920
6.38



695
963
4.96



617
824
2.84


Alloy 48
997
1303
4.17



1082
1390
2.27



1071
1369
3.40


Alloy 49
1018
1252
3.92



1049
1151
2.47



1047
1133
2.13


Alloy 50
904
991
1.22



1024
1074
1.27



981
1127
2.02


Alloy 51
624
892
5.39



599
846
4.67



613
911
6.56








Alloy 52
Not tested (brittle)










Alloy 53
946
1265
4.49



937
1130
2.79



851
1251
4.80


Alloy 54
1077
1218
1.77



1142
1386
2.57



1098
1244
1.98


Alloy 55
915
1172
4.07



869
1147
5.90



938
1200
4.57


Alloy 56
998
998
1.22


Alloy 57
804
1123
3.13



688
1038
5.13



686
862
2.07


Alloy 58
1001
1298
1.70


Alloy 59
1159
1627
4.67



1260
1638
2.35



1391
1512
1.92


Alloy 60
695
888
0.88









Applications
The alloys herein in either forms as Class 1 or Class 2 Steel have a variety of applications. These include but are not limited to structural components in vehicles, including but not limited to parts and components in the vehicular frame, front end structures, floor panels, body side interior, body side outer, rear structures, as well as roof and side rails. While not all encompassing, specific parts and components would include B-pillar major reinforcement, B-pillar belt reinforcement, front rails, rear rails, front roof header, rear roof header, A-pillar, roof rail, C-pillar, roof panel inners, and roof bow. The Class 1 and/or Class 2 steel will therefore be particular useful in optimizing crash worthiness management in vehicular design and allow for optimization of key energy management zones, including engine compartment, passenger and/or trunk regions where the strength and ductility of the disclosed steels will be particular advantageous.
The alloys herein may also provide for use in additional non-vehicular applications, such as for drilling applications, which therefore may include use as a drill collars (a component that provides weight on a bit for drilling), drill pipe (hollow wall pipe used on drilling rigs to facilitate drilling), tool joints (i.e. the threaded ends of drill pipe) and wellheads (i.e. the component of a surface or an oil or gas well that provides the structural and pressure-containing interface for drilling and production equipment) including but not limited to ultra-deep and ultra-deep water and extended reach (ERD) well exploration.

## Claims

1. A method comprising:
supplying a metal alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.50 atomic percent, B at 4.0 to 8.0 atomic percent, Si at 4.0 to 8.0 atomic percent;
melting said alloy and solidifying to provide a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm;
mechanical stressing said alloy and/or heating to form at least one of the following grain size distributions and mechanical property profiles, wherein said boride grains provide pinning phases that resist coarsening of said matrix grains:
(a) matrix grain size in the range of 500 nm to 20,000 nm, boride grain size in the range of 25 nm to 500 nm, precipitation grain size in the range of 1 nm to 200 nm wherein said alloy indicates a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa and tensile elongation of 10 to 40%; or
(b) matrix grain size in the range of 100 nm to 2000 nm and boride grain size in the range of 25 nm to 500 nm which has a yield strength of 300 MPa to 600 MPa.

2. The method of claim 1 wherein said alloy includes one or more of the following:
V at 1.0 to 3.0 atomic percent;
Zr at 1.0 atomic percent;
C at 0.2 to 3.0 atomic percent;
W at 1.0 atomic percent; or
Mn at 0.2 to 4.6 atomic percent.

3. The method of claim 1 wherein said melting is achieved at temperatures in the range of 1100° C. to 2000° C. and solidification is achieved by cooling in the range of 11×103 to 4×10−2 K/s.

4. The method of claim 1 wherein said alloy having said grain size distribution (b) is exposed to a stress that exceeds said yield strength of 300 MPa to 600 MPa wherein said grain size remains in the range of 100 nm to 2000 nm, said boride grain size remains in the range of 25 nm to 500 nm, along with the formation of precipitation grains of 1 nm to 200 nm wherein said precipitation grains include a hexagonal phase.

5. The method of claim 4 wherein said alloy indicates a tensile strength of 720 MPa to 1580 MPa and an elongation of 5% to 35%.

6. The method of claim 5 wherein said alloy indicates a strain hardening coefficient of 0.2 to 1.0.

7. The method of claim 1 wherein said alloy having said mechanical property profile and grain size distribution (a) or (b) is in the form of sheet.

8. The method of claim 4 wherein said alloy having said grain size in the range of 100 nm to 2000 nm, said boride grain size in the range of 25 nm to 500 nm, and said precipitation grains in the range of 1 nm to 200 nm wherein said precipitation grains include a hexagonal phase, is in the form of sheet.

9. The method of claim 1 wherein said alloy having said mechanical property profile and grain size distribution (a) is positioned in a vehicle.

10. The method of claim 5 wherein said alloy is positioned in a vehicle.

11. The method of claim 1 wherein said alloy having said mechanical property profile and grain size distribution is positioned in one of a drill collar, drill pipe, tool joint or wellhead.

12. The method of claim 5 wherein said alloy is positioned in one of a drill collar, drill pipe, tool joint or wellhead.

13. A method comprising:
supplying a metal alloy comprising Fe at a level of 53.5 to 72.1 atomic percent, Cr at 10.0 to 21.0 atomic percent, Ni at 2.8 to 14.50 atomic percent, B at 4.00 to 8.00 atomic percent, Si at 4.00 to 8.00 atomic percent;
melting said alloy and solidifying to provide a matrix grain size in the range of 500 nm to 20,000 nm containing 10% to 70% by volume ferrite and a boride grain size in the range of 25 nm to 500 nm wherein said boride grains provide pinning phases that resist coarsening of said matrix grains upon application of heat and wherein said alloy has a yield strength of 300 MPa to 600 MPa;
heating said alloy wherein said grain size is in the range of 100 nm to 2000 nm, said boride grain size remains in the range of 25 nm to 500 nm and said level of ferrite increases to 20% to 80% by volume;
stressing said alloy to a level that exceeds said yield strength of 300 MPa to 600 MPa wherein said grain size remains in the range at 100 nm to 2000 nm, said boride grain size remains in the range of 25 nm to 500 nm, along with the formation of precipitation grains in the range of 1 nm to 200 nm and said alloy has a tensile strength of 720 MPa to 1580 MPa and an elongation of 5% to 35%.

14. The method of claim 13 wherein said alloy includes one or more of the following:
V at 1.0 to 3.0 atomic percent;
Zr at 1.0 atomic percent;
C at 0.2 to 3.0 atomic percent;
W at 1.00 atomic percent; or
Mn at 0.20 to 4.6 atomic percent.

15. The method of claim 13 wherein said melting is achieved at temperature in the range of 1100° C. to 2000° C. and solidification is achieved by cooling in the range of 11×103 to 4×10−2K/s.

16. The method of claim 13 wherein said alloy is in the form of sheet.

17. A metallic alloy comprising:
Fe at a level of 53.5 to 72.1 atomic percent;
Cr at 10.0 to 21.0 atomic percent;
Ni at 2.8 to 14.5 atomic percent;
B at 4.0 to 8.0 atomic percent;
Si at 4.0 to 8.0 atomic percent;
wherein said alloy indicates a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm and wherein said alloy having been exposed to mechanical stress and/or heat to indicate at least one of the following:
(a) exposure to mechanical stress said alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa, tensile elongation of 10 to 40%; or
(b) exposure to heat, followed by mechanical stress, said alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 1300 MPa, tensile strength of 720 MPa to 1580 MPa, tensile elongation of 5.0% to 35.0%.

18. The metallic alloy of claim 17 wherein said mechanical property profile (a) includes a strain hardening coefficient of 0.1 to 0.4.

19. The metallic alloy of claim 17 wherein said mechanical property profile (b) includes a strain hardening coefficient of 0.2 to 1.0.

20. The metallic alloy of claim 17 wherein said mechanical property profile (a) comprises the following grain size distribution: a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm and a precipitation grain size in the range of 1.0 nm to 200 nm.

21. The metallic alloy of claim 17 wherein said mechanical property profile (b) comprise the following grain size distribution: a matrix grain size in the range of 100 nm to 2000 nm, a boride grain size in the range of 25 nm to 500 nm and precipitation grain size in the range of 1 nm to 200 nm.

22. The metallic alloy of claim 21 wherein said precipitation grain size of 1 nm to 200 nm includes a hexagonal phase.

23. The metallic alloy of claim 17 wherein said alloy includes one or more of the following:
V at 1.0 to 3.0 atomic percent;
Zr at 1.0 atomic percent;
C at 0.2 to 3.0 atomic percent;
W at 1.0 atomic percent; or
Mn at 0.2 to 4.6 atomic percent.

24. The alloy of claim 17 wherein said alloy recited in (a) or (b) is in the form of sheet material.

25. A metallic alloy comprising:
Fe at a level of 53.5 to 72.1 atomic percent;
Cr at 10.0 to 21.0 atomic percent;
Ni at 2.8 to 14.5 atomic percent;
B at 4.0 to 8.0 atomic percent;
Si at 4.0 to 8.0 atomic percent;
wherein said alloy indicates a matrix grain size in the range of 500 nm to 20,000 nm and a boride grain size in the range of 25 nm to 500 nm and wherein said alloy having been exposed to mechanical stress and/or heat to indicate at least one of the following:
(a) exposure to mechanical stress said alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 840 MPa, tensile strength of 630 MPa to 1100 MPa, tensile elongation of 10% to 40%, and a matrix grain size in the range of 500 nm to 20,000 nm, a boride grain size in the range of 25 nm to 500 nm and a precipitation grain size in the range of 1.0 nm to 200 nm; or
(b) exposure to heat followed by mechanical stress, said alloy indicates a mechanical property profile providing a yield strength of 300 MPa to 1300 MPa, tensile strength of 720 MPa to 1580 MPa, tensile elongation of 5% to 35% and a matrix grain size in the range of 100 nm to 2000 nm, a boride grain size in the range of 25 nm to 500 nm, and a precipitation grain size in the range of 1 nm to 200 nm.

26. The metallic alloy of claim 25 wherein said alloy includes one or more of the following:
V at 1.0 to 3.0 atomic percent;
Zr at 1.0 atomic percent;
C at 0.2 to 3.00 atomic percent;
W at 1.0 atomic percent; or
Mn at 0.20 to 4.6 atomic percent.

27. The alloy of claim 17 wherein said mechanical property profile (a) includes a strain hardening coefficient of 0.1 to 0.4 and said mechanical property profile (b) includes a strain hardening coefficient of 0.2 to 1.0.