Patent Application
20130233452
This application deals with new class of non-stainless steel alloys with advanced property combination applicable to sheet production by methods such as chill surface processing.
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 cooling rate of the steel at solidification or thermal treatment, 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.
Non-stainless steels may be understood herein to contain less than 10.5% of chromium and are typically represented by plain carbon steel which is by far the most widely used kind of steel. The properties of carbon steel depend primarily on the amount of carbon it contains. With very low carbon content (below 0.05% C), these steels are relatively ductile and have properties similar to pure iron. They cannot be modified by heat treatment. They are inexpensive, but engineering applications may be restricted to non-critical components and general paneling work.
Pearlite structure formation in most alloy steels requires less carbon than in ordinary carbon steels. The majority of these alloy steels is low carbon material and alloyed with a variety of elements in total amounts of between 1.0% and 50% by weight to improve its mechanical properties. Lowering the carbon content to the range of 0.10% to 0.30%, along with some reduction in alloying elements increases the weldability and formability of the steel while maintaining its strength. Such alloys are classed as a high-strength low-alloy steels (HSLA) exhibiting tensile strengths from 270 to 700 MPa.
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, low-strength steel (LSS), high-strength steel (HSS) and AHSS may indicate tensile elongations at levels of 25%-55%, 10%-45% and 4%-30%, respectively.
Much higher strength (up to 2500 MPa) has been achieved in maraging steels which are carbon free iron-nickel alloys with additions of cobalt, molybdenum, titanium and aluminum. The term maraging is derived from the strengthening mechanism, which is transforming the alloy to martensite with subsequent age hardening. The common, non stainless grades of maraging steels contain 17% to 18% nickel, 8% to 12% cobalt, 3% to 5% molybdenum and 0.2% to 1.6% titanium. The relatively high price of maraging steels (they are several times more expensive than the high alloy tool steels produced by standard methods) significantly restricts their application in many areas (for example, automotive industry). They are highly sensitive to nonmetallic inclusions, which act as stress raisers and promote nucleation of voids and microcracks leading to a decrease in ductility and fracture toughness of the steel. To minimize the content of nonmetallic inclusions, the maraging steels are typically melted under vacuum resulting in high cost processing.
The present disclosure relates to a method for producing a metallic alloy comprising a method comprising supplying a metal alloy comprising Fe at a level of 65.5 to 80.9 atomic percent, Ni at 1.7 to 15.1 atomic percent, B at 3.5 to 5.9 atomic percent, Si at 4.4 to 8.6 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 and a boride grain size of 25 nm to 500 nm. One may then mechanical stress said 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 of 500 nm to 20,000 nm, boride grain size of 25 nm to 500 nm, precipitation grain size 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) refined matrix grain size of 100 nm to 2000 nm, precipitation grain size of 1 nm to 200 nm, boride grain size of 200 nm to 2,500 nm where the alloy has a yield strength of 300 MPa to 600 MPa. The alloy having the refined grain size distribution (b) may be exposed to a stress that exceeds the yield strength of 300 MPa to 600 MPa wherein the refined grain size remains at 100 nm to 2000 nm, the boride grain size remains at 200 nm to 2500 nm, the precipitation grains remain at 1 nm to 200 nm, wherein said alloy indicates a yield strength of 300 MPa to 1400 MPa, tensile strength of 875 MPa to 1590 MPa and an elongation of 5% to 30%.
The present disclosure also relates to a method comprising supplying a metal alloy comprising Fe at a level of 65.5 to 80.9 atomic percent, Ni at 1.7 to 15.1 atomic percent, B at 3.5 to 5.9 atomic percent, Si at 4.4 to 8.6 atomic percent. One may then melt the alloy and solidify to provide a matrix grain size of 500 nm to 20,000 nm and a boride grain size of 100 nm to 2500 nm. This may then be followed by heating the alloy and forming lath structure including grains of 100 nm to 10,000 nm and boride grain size of 100 nm to 2500 nm wherein the alloy has a yield strength of 300 MPa to 1400 MPa, tensile strength of 350 MPa to 1600 MPa and elongation of 0-12%. One may then heat the aforementioned lath structure and form lamellae grains 100 nm to 10,000 nm thick, 0.1-5.0 microns in length and 100 nm to 1000 nm in width along with boride grains of 100 nm to 2500 nm and precipitation grains of 1 nm to 100 nm, wherein the alloy indicates a yield strength of 350 MPa to 1400 MPa. The aforementioned lamellae structure may undergo a stress and form an alloy having grains of 100 nm to 5000 nm, boride grains of 100 nm to 2500 nm, precipitation grains of 1 nm to 100 nm where the alloy has a yield strength of 350 MPa to 1400 MPa, a tensile strength of 1000 MPa to 1750 MPa and elongation of 0.5% to 15.0%.
The present disclosure further relates to metallic alloy comprising Fe at a level of 65.5 to 80.9 atomic percent; Ni at 1.7 to 15.1 atomic percent; B at 3.5 to 5.9 atomic percent; and Si at 4.4 to 8.6 atomic percent, wherein the alloy indicates a matrix grain size of 500 nm to 20,000 nm and boride grain size of 100 nm to 2500 nm. The alloy, upon a first exposure to heat forms a lath structure including grains of 100 nm to 10,000 nm and boride grain size of 100 nm to 2500 nm wherein the alloy has a yield strength of 400 MPa to 1400 MPa, tensile strength of 350 MPa to 1600 MPa and elongation of 0-12%. Upon a second exposure to heat followed by stress the alloy has grains of 100 nm to 5000 nm, boride grains of 100 nm to 2500 nm, precipitation grains of 1 nm to 100 nm and the alloy has a yield strength of 350 MPa to 1400 MPa, a tensile strength of 1000 MPa to 1750 MPa and elongation of 0.5% to 15.0%.
The detailed description below may be better understood with reference to the accompanying figures which are provided for illustrative purposes and not to be considered as limiting any aspect of this invention.
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 with widths in the range of 100 to 5000 mm. These thickness ranges and width ranges may be adjusted in these ranges at 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 4×10−2 K/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, centrifugal casting etc. Also, powder metallurgy through either conventional press and sintering or through HIPing/forging is a contemplated route to make partially or fully dense parts and devices utilizing the chemistries, structures, and mechanisms described in this application (i.e. the Class 2 or Class 3 Steel described herein).
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
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
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.
The non-stainless steel alloys herein are such that they are capable of formation of what is described herein as Class 1, Class 2 Steel or Class 3 Steel which are preferably crystalline (non-glassy) with identifiable crystalline grain size morphology. The ability of the alloys to form Class 2 or Class 3 Steels herein is described in detail herein. However, it is useful to first consider a description of the general features of Class 1, Class 2 and Class 3 Steels, which is now provided below.
Class 1 Steel
The formation of Class 1 Steel herein (non-stainless) is illustrated in
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
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=K εn, where σ represents the applied stress on the material, ε 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.
Class 2 Steel
The formation of Class 2 Steel herein (non-stainless) is illustrated in
As shown therein, Structure #1 is initially formed in 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. Grain size herein may again 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 2 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 2 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 are referenced 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 and which are unaffected by Mechanisms #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
Characteristic of the Static Nanophase Refinement Mechanism #1 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 (alpha-iron) initially present in the modal structure (Structure 1) of Class 2 steel is 0 to 45%. The volume fraction of ferrite (alpha-iron) in Structure #2 as a result of Static Nanophase Refinement Mechanism #2 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 rather than 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 875 to 1590 MPa with 5 to 30% total elongation.
Depending on alloy chemistries, nano-scale precipitates can form during Static Nanophase Refinement and the subsequent thermal process in some of the non-stainless high-strength steels. The nano-precipitates are in the range of 1 nm to 200 nm, with the majority (>50%) of these phases 10˜20 nm in size, which are much smaller than the boride pinning phase formed in Structure #1 for retarding matrix grain coarsening. Also, during Static Nanophase Refinement, the boride grain sizes grow larger to a range from 200 to 2500 nm in size.
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 (
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 but only 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 point of the material. The tensile properties that can be achieved for alloys that achieve Structure 3 include tensile strength values in the range from 875 to 1590 MPa and 5 to 30% total elongation. The level of tensile properties achieved is also dependent on the amount of transformation occurring as the strain increases corresponding to the characteristic stress strain curve for a Class 2 steel.
Thus, depending on the level of transformation, 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 1400 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 1400 MPa) as applied to Structure #2 transformation into Structure #3, allowing tunable variations to enable both the designer and end users in a variety of applications, and utilize Structure #3 in various applications such as crash management in automobile body structures.
With regards to this dynamic mechanism shown in
Note that dynamic recrystallization is a known process but differs from Mechanism #2 (
Class 3 Steel
Class 3 steel (non-stainless) is associated with formation of a High Strength Lamellae NanoModal Structure through a multi-step process as now described herein.
In order to achieve a tensile response involving high strength with adequate ductility in non-stainless carbon-free steel alloys, a preferred seven-step process is now disclosed and shown in
Structure #1 involving a formation of the Modal Structures (i.e. bi, tri, and higher order) may be achieved in the alloys with the referenced chemistries in this application by processing through the laboratory scale as shown and/or through industrial scale methods involving chill surface processing such as twin roll casting or thin slab casting. The Modal Structure of Class 3 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 ferrite or alpha-Fe (required) and optionally austenite or gamma-Fe; and (2) boride grain size of 100 nm to 2500 nm (i.e. non-metallic grains such as M2B where M is the metal and is covalently bonded to B); (3) yield strengths of 350 to 1000 MPa; (4) tensile strengths of 200 to 1200 MPa; and total elongation of 0-3.0%. It will also indicate dendritic growth morphology of the matrix grains. 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 and which are unaffected by Mechanism #1, #2 or #3 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. Accordingly, Structure #1 of Class 3 steel herein includes ferrite along with such boride phases.
Structure #2 involves the formation of the Modal Lath Phase Structure with uniformly distributed precipitates from Modal Structure (Structure 1) with dendritic morphology though Mechanism #1. Lath phase structure may be generally understood as a structure composed from plate-shaped crystal grains. Reference to “dendritic morphology” may be understood as tree-like and reference to “plate shaped” may be understood as sheet like. Lath structure formation preferably occurs at elevated temperature (e.g. at temperatures of 700° C. to 1200° C.) through plate-like crystal grain formation with: (1) lath structural grain sizes typically from 100 to 10,000 nm; (2) boride grain size of 100 nm to 2,500 nm; (3) yield strengths of 300 MPa to 1400 MPa; (4) tensile strengths of 350 MPa to 1600 MPa; (5) elongation of 0-12%. Structure #2 also contains alpha-Fe and gamma-Fe remains optional.
A second phase of boride precipitates with a size typically from 100 to 1000 nm may be found distributed in the lath matrix as isolated particles. The second phase of boride precipitates may be understood as non-metallic grains of different stoichiometry (M2B, M3B, MB (M1B1), M23B6, and M7B3) where M is the metal and is covalently bonded to Boron. These boride precipitates are distinguished from the boride grains in Structure #1 with little or no change in size.
Structure #3 (Lamellae NanoModal Structure) involves the formation of the lamellae morphology as a result of static transformation of ferrite into one or several phases through Mechanism #2 identified as Lamellae Nanophase Creation. Static transformation is a decomposition of the parent phase into new phase or several new phases due to alloying elements distribution by diffusion during elevated temperature heat treatment, which may preferably occur in the temperature range from 700° C. to 1200° C. Lamellae (or layered) structure is composed of alternating layers of two phases whereby individual lamellae exist within a colony connected in three dimensions. A schematic illustration of lamellae structure is shown in
In Class 3 alloys, Lamellae Nanomodal Structure contains: (1) lamellas of 100 nm to 1000 nm wide with a thickness in the range of 100 nm to 10,000 nm with a length of 0.1 to 5 microns; (2) boride grains of 100 nm to 2500 nm of different stoichiometry (M2B, M3B, MB (M1B1), M23B6, and M7B3) where M is the metal and is covalently bonded to Boron, (3) precipitation grains of 1 nm to 100 nm; (4) yield strength of 350 MPa to 1400 MPa. The Lamellae Nanomodal Structure continues to contain alpha-Fe and gamma-Fe remains optional.
Lamellae NanoModal Structure (Structure #3) transforms into Structure #4 through Dynamic Nanophase Strengthening (Mechanism #3, exposure to mechanical stress) during plastic deformation (i.e. exceeding the yield stress for the material) displaying relatively high tensile strengths in the range of 1000 MPa to 1750 MPa. In
The strengthening during deformation is related to phase transformation that occurs as the material strains under stress and defines Mechanism #3 as a dynamic process. For the alloy to display high strength at the level described in this application, lamellae structure is preferably formed prior to deformation. Specific to this mechanism, the micron scale austenite phase is transformed into new phases with reductions in microstructural feature scales generally down to the nanoscale regime. Some fraction of austenite may initially form in some Class 3 alloys during casting and then may remain present in Structure #1 and Structure #2. During straining when stress is applied, new or additional phases are formed with nanograins typically in a range from 1 to 100 nm. See Table 15.
In the post-deformed Structure #4 (High Strength Lamellae NanoModal Structure), the ferrite grains contain alternating layers with nanostructure composed from new phases formed during deformation. Depending on the specific chemistry and the stability of the austenite, some austenite may be additionally present. In contrast with layers in Structure #3 where each layer represents a single or just few grains, in Structure #4, a large number of nanograins of different phases are present as a result of Dynamic NanoPhase Strengthening. Since nanoscale phase formation occurs during alloy deformation, it represents a stress induced transformation and defined as a dynamic process. Nanoscale phase precipitations during deformation are responsible for extensive strain hardening of the alloys.
The dynamic transformation can occur partially or completely and results in the formation of a microstructure with novel nanoscale/near nanoscale phases specified as High Strength Lamellae NanoModal Structure (Structure #4) that provides high strength in the material. Thus the Structure #4 can be formed with various levels of strengthening depending on specific chemistry and the amount of strengthening achieved by Mechanism #3. Table 2 below provides a comparison of the structure and performance features of Class 3 Steel herein.
The formation of Modal Structure (MS) in either Class 2 or Class 3 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 4×10−2 K/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 subjected to heating at a 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 and/or NMS 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. The amount of DNS that occurs may depend on the volume fraction of Static Nanophase Refinement in the material prior to 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
With respect to Class 3 Steel herein, Mechanism #1 which is the Lath Phase Creation occurs during elevated temperature exposure of the initial Modal Structure #1 and can occur during Stage 1, Stage 2 or Stage 3 (after MS formation) of twin roll production or thin slab casting production. In some alloys, Lath Structure Creation can occur at solidification at Stage 1 of twin roll or thin slab casting production. Mechanism #1 results in formation of Modal Lath Phase Structure specified as Structure #2. The formation of Structure #2 is critical step in terms of further Lamellae NanoModal Structure (Structure #3) formation through Mechanism #2 specified as Lamellae Nanophase Creation by phase transformation. Mechanism #2 in the sheet alloys can occur during Stage 1, 2, or 3 of twin roll production or thin slab casting production or during post processing of the sheets. In some alloys, Structure #3 may also form at earlier Stages of casting production such as Stage 2 or Stage 3 of twin roll production or thin slab casting, as well as at post-processing treatment of produced sheet. Lamellae NanoModal Structure is responsible for high strength of the alloys of current application and has ability for strengthening during room temperature deformation through Mechanism #3 specified as Dynamic Nanophase Strengthening. The level of Dynamic Nanophase Strengthening that occurs will depend on the alloy chemistry and on a stress level induced into the sheet. The strengthening may also occur during subsequent post processing of sheets produced by twin roll production or thin slab casting into final parts involving hot or cold forming of the sheets. Thus, the resultant High Strength Lamellae NanoModal Structure specified as Structure #4 can occur at post-processing of produced sheets by methods that involve mechanical deformation to different levels of strengthening depending on the alloy chemistry, deformation parameters and post-deformation thermal cycle(s).
Preferred Alloy Chemistries and Sample Preparation The chemical composition of the alloys studied is shown in Table 3 which provides the preferred atomic ratios utilized. These chemistries have been used for material processing through plate 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 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plates with thickness of 1.8 mm mimicking alloy solidification into a sheet with similar thickness between rolls at Stage 1 of Twin Roll Casting process.
Accordingly, in the broad context of the present disclosure, the alloy chemistries that may preferably be suitable for the formation of the Class 1, Class 2 or Class 3 Steel herein, include the following whose atomic ratios add up to 100. That is, the alloys may include Fe, Ni, B and Si. The alloys may optionally include Cr, Cu and/or Mn. Preferably, with respect to atomic ratios, the alloys may contain Fe at 65.64 to 80.85, Ni at 1.75 to 15.05, B at 3.50 to 5.82 and Si at 4.40 to 8.60. Optionally, and again in atomic ratios, one may also include Cr at 0 to 8.72, Cu at 0 to 2.00 and Mn at 0-18.74. Accordingly, the levels of the particular elements may be adjusted to 100 as noted above. Impurities known/expected to be present include, but are not limited to, C, Al, Mo, Nb, Ti, S, O, N, P, W, Co, and Sn. Such impurities may be present at levels up to 10 atomic percent.
The atomic ratio of Fe present may therefore be 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, 69.1, 69.2, 69.3, 69.4, 69.5, 69.6, 69.7, 69.8, 69.9, 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, 72.2, 72.3, 72.4, 72.5, 72.6, 72.7, 72.8, 72.9, 80.0, 80.1, 80.2, 80.3, 80.4, 80.5, 80.6, 80.7, 80.8, 80.9. The atomic ratio of Ni may therefore be 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, 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, 14.5, 14.6, 14.7, 14.8, 14.9, 15.0, 15.1. The atomic ratio of B may therefore be 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. The atomic ratio of Si may therefore be 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.
The atomic ratios of the optional elements such as Cr may therefore be 0.1, 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, 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., 7.6, 7.7, 7.8, 7.9, 8.0, 8.1, 8.2, 8.3, 8.4, 8.5, 8.6, 8.7, and 8.8. The atomic ratio of Cu if present may therefore be 0.1, 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 and 2.0. The atomic ratio of Mn if present may therefore be 0.1, 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, 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, 14.5, 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 and 18.8.
The alloys may herein also be more broadly described as an Fe based alloy (greater than 50.00 atomic percent) and including B, Ni and Si and capable of forming the indicated structures (Class 1, Class 2 and/or Class 3 Steel) and/or undergoing the indicated transformations upon exposure to mechanical stress and/or mechanical stress in the presence of heat treatment/thermal exposure. 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.
Thermal analysis was done on the as-solidified cast plate samples on a NETZSCH DSC 404F3 PEGASUS V5 system. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) were performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultra-high purity argon. In Table 4, elevated temperature DTA results are shown indicating the melting behavior for the alloys shown in Table 3. As can be seen from the tabulated results in Table 4, the melting occurs in 1, 2, 3 or 4 stages with initial melting observed from ˜1108° C. depending on alloy chemistry. Final melting temperature is up to ˜1400° C. Variations in melting behavior may also reflect complex phase formation at chill surface processing of the alloys depending on their chemistry.
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 5 and was found to vary from 7.48 g/cm3 to 7.71 g/cm3. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3.
The tensile specimens were cut from selected plates 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. Video extensometer was utilized for strain measurements. In Table 6, a summary of the tensile test results including total tensile elongation (strain), yield stress, and ultimate strength are listed for selected as-cast plates. The mechanical characteristic values strongly depend on alloy chemistry and processing condition as will be showed later. As can be seen, the tensile strength values in these selected alloys vary from 350 to 1196 MPa. The total elongation value varied from 0.22 to 2.80% indicating limited ductility of alloys in as-cast state. In some specimens, failure occurred in elastic region at stress as low as 200 MPa and yielding was not reached.
Properties in Table 6 are related to the formation of the Structure #1 (
Each plate 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 plates 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 7. The key aspect of the HIP cycle was to remove macrodefects such as pores and small inclusions 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 of a plate before and after HIP cycle is shown in
The tensile specimens were cut from the plates after HIP cycle 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 8, a summary of the tensile test results including total tensile elongation (strain), yield stress, and ultimate tensile strength are shown for the cast plates after HIP cycle. Additional column is added that specifies the alloy mechanical response in correspondence with the class of behavior (
Properties of the alloys that demonstrated Class 3 behavior in Table 8 are related to the formation of the Structure #2 (
Properties of the alloys that demonstrated Class 2 behavior in Table 8 are related to the formation of the Structure #2 (
After HIP cycle, the plate material was heat treated in a box furnace at parameters specified in Table 9. The 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. 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.
The tensile specimens were cut from the plates 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 10, a summary of the tensile test results including total tensile elongation (strain), yield stress, and ultimate tensile strength are shown for the cast plates after HIP cycle and heat treatment. Additional column is added that specifies the alloy mechanical response in correspondence with the class of behavior (
In the case of Class 2 behavior, the tensile strength of the alloys (Structure 3 in Table 2) varies from 875 to 1590 MPa. The total elongation value varies from 5.0 to 30.0% providing superior high strength/high ductility property combination. Such property combination related to the formation of the Structure #3 (
In a case of Class 3 behavior, the tensile strength of the alloys is equal to or higher than 1000 MPa and the data varies from 1004 to 1749 MPa. The total elongation values for the sample alloys vary from 0.5 to 14.5%. High strength of the alloys in Table 10 with Class 3 behavior related to the formation of Structure #3 (
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 11. Tensile stress-strain curves are compared to that of existing Dual Phase (DP) steels (
According to the alloy stoichiometries in Table 3, the Alloy 51 was weighed out using high purity elemental charges. It should be noted that Alloy 51 has demonstrated Class 2 behavior with high tensile ductility at high strength. The resulting charges were arc-melted into several (usually 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 plates under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. Two of the plates were then HIPed at 1100° C. for 1 hour. One of the HIPed plates was then subsequently heat treated at 700° C. for 1 hour with air cooling to room temperature. The plates 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 51 plates 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 51 plate sample in the as-cast, HIPed and HIPed/heat treated conditions are shown in
Additional details of the Alloy 51 plate 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 45 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
In the as-cast plate, two phases were identified, cubic γ-Fe (austenite) and a complex mixed transitional metal boride phase with the M2B1 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 would exhibit a lattice parameter equal to a=3.575 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 it is likely that silicon is also dissolved into this structure so it is not a pure boride phase. Additionally, as can be seen in Table 12, while the phases do not change, the lattice parameters do change as a function of the plate condition (i.e. as-cast, HIPed, HIPed/heat treated), which indicates that redistribution of alloying elements is occurring.
As can be seen in Table 12, after the HIP exposure (1100° C. for 1 hour at 15 ksi) three phases are found which are α-Fe (ferrite), M2B1 phase, and γ-Fe (austenite). Note that α-Fe is believed to be formed from the γ-Fe (austenite) phase. Note also that the lattice parameters of the M2B1 and γ-Fe phases are different indicating that elemental redistribution/diffusion is occurring. As can be seen in Table 12, after the heat treatment at 700° C. for 1 hour, four phases are present which are α-Fe (ferrite), M2B1 phase, and two newly identified hexagonal phases. Note that γ-Fe is not found in the sample after heat treatment indicating that this phase transformed into the newly found phases. The M2B1 phase is still present in the X-ray diffraction scan but its lattice parameters have changed significantly indicating that atomic diffusion has occurred at elevated temperature. One identified new hexagonal phase is representative of a ditrigonal dipyramidal class and has a hexagonal P6bar2C space group (#190) and the other newly identified hexagonal phase is representative of a dihexagonal pyramidal class and has a hexagonal P63mc space group (#186). It is theorized based on the small crystal unit cell size that the ditrigonal dipyramidal phase is likely a silicon based phase possibly a previously unknown Si—B phase which may be stabilized by the presence of the additional alloying elements in the stoichiometry. Also note that based on the ratio of peak intensities it appears that the dihexagonal pyramidal may be forming with specific orientation relationships since the diffracted intensity from the (002) planes is much higher than expected and the diffracted intensity from the (103) and (112) planes is much lower. Based on the ratio of peak intensities, it seems that one of the major differences of the heat treatment is the creation of a lot more of the ditrigonal dipyramidal hexagonal phase.
To examine the structural features of the Alloy 51 plates in more detail, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM samples, specimens were cut from the as-cast, HIPed, and HIPed/heat-treated plates, and then ground and polished to a thickness of ˜30 to ˜40 μm. Discs of 3 mm in diameter were then punched from these polished thin samples, 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
According to the alloy stoichiometries in Table 3, the Alloy 6 that represents Class 3 alloy was weighed out from high purity elemental charges. It should be noted that Alloy 6 has demonstrated Class 3 behavior with very high strength characteristics. 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 plates under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. Two of the plates were then HIPed at 1100° C. for 1 hour. One of the HIPed plates was then subsequently heat treated at 700° C. for 1 hour with slow cooling to room temperature (670 minutes total time). The plates in the as-cast, HIPed and HIPed/heat treated states were then cut by 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 6 plates 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 manufactured by Carl Zeiss SMT Inc. Example SEM backscattered electron micrographs of the plate microstructure in the as-cast, HIPed and HIPed and heat treated conditions are shown in
Similar to Class 2 alloy, in the as-cast sample from Class 3 alloy, the microstructure contains two basic components, i.e., the matrix dendrite grains and an intergranular area, as marked by A and B in
Additional details of the Alloy 6 plate 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 45 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
In the as-cast plate and HIPed (1100° C. for 1 hour) plate, two phases were identified, cubic α-Fe (ferrite) and a complex mixed transitional metal boride phase with the M2B1 stoichiometry. Note that the lattice parameters of the identified phases are different from that found for pure phases clearly indicating the dissolution of the alloying elements. For example, 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. This is consistent with the SEM studies which did not show new phases present but homogenization of the structure. After the heat treatment (700° C. slow cool to room temperature (670 minute total time)) as can be seen in Table 13, the α-Fe (ferrite) and M2B1 phases are all present although the lattice parameters change indicating diffusion and redistribution of the alloying elements. Additionally, γ-Fe (not a pure phase since it exhibits a lattice parameter of a=3.577 Å which is slightly larger than that of a pure phase at (a=3.575 Å)) and a newly identified hexagonal phase is representative of a dihexagonal pyramidal class and has a hexagonal P63mc space group (#186) are found in the X-ray diffraction pattern. The presence of these new phases is consistent with the new precipitates found in the SEM studies and contributes to the formation of the lath matrix structure.
To examine the structural details of the Alloy 6 plates 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 plates. 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 solution. The prepared specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope (TEM) operated at 200 kV.
TEM analysis was conducted at both the intergranular region and the matrix grains. As shown in
During heat treatment, the boride precipitates grow slightly, but the lath structure in the matrix experiences great changes.
The tensile properties of the steel plate produced in this application will be sensitive to the specific structure and specific processing conditions that the plate experiences. In
Samples that were cut out of the Alloy 51 tensile gage and grip section 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 manufactured by Carl Zeiss SMT Inc. Example SEM backscattered electron micrographs from tensile gage section and grip section are shown in
For the Alloy 51 plate HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 1 hour with air cooling, additional structural details were obtained through using X-ray diffraction which was done on both the undeformed plate 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 45 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
The X-ray pattern for the deformed Alloy 51 tensile tested specimen (HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 1 hour with air cooling) was subsequently analyzed using Rietveld analysis using Siroquant software. As shown in
To examine the structural changes of the Alloy 51 plates 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 samples, 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
The very fine precipitates observed by TEM would include the new hexagonal phases produced by heat treatment and by deformation, identified by X-ray diffraction (see section above). Due to the pinning effect by the precipitates, the matrix grains are refined to a higher level thanks to the dislocation accumulation that increases the grain lattice misorientation during the tensile deformation. While the deformation-induced nanoscale phase formation may contribute to the hardening in the Alloy 51 plate, the work-hardening of Alloy 51 is strengthened by dislocation based mechanisms including dislocation pinning by precipitates.
As it was shown, the Alloy 51 plate has demonstrated Structure #1 Modal Structure (Step #1) in as-cast state (
The tensile properties of the steel plate produced in this application will be sensitive to the specific structure and specific processing conditions that the plate experiences. In
For the Alloy 6 plate HIPed at 1100° C. for 1 hour, additional structural details were obtained through using X-ray diffraction which was done on both the undeformed plate 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 45 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
The X-ray pattern for the deformed Alloy 6 tensile tested specimen (HIPed (1100° C. for 1 hour) was subsequently analyzed using Rietveld analysis using Siroquant software. As shown in
To focus on structural changes occurring during tensile testing, the Alloy 6 plate HIPed at 1100° C. for 1 hour, and heat treated at 700° C. for 60 minutes with slow furnace cooling was examined by TEM. TEM specimens were prepared from HIPed and heat treated plate both in the undeformed state and after tensile testing until failure. TEM specimens were made from the plate 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 plate specimens. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV.
Using high purity elements, 35 g alloy feedstocks of the Alloy 17 and Alloy 27 were weighed out according to the atomic ratios provided in Table 3. The only difference between these two alloys is that ½ of Ni in Alloy 17 is substituted by Mn in Alloy 27. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 from the Alloy 17 and Alloy 27 were subjected to a HIP cycle C (at 1100° 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 plates 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 HIP cycle, the plates were heat treated at 700° C. for 1 h with air cooling. Tensile specimens were cut from the treated plates.
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. Representative curves for both alloys are shown in
Samples from both alloys after tensile testing were examined by SEM. Samples were cut from the gage section and then 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 manufactured by Carl Zeiss SMT Inc. SEM backscattered images of the sample microstructure are shown in
In the Alloy 17 sample, the dark boride pinning phase (mostly 1˜2 μm in diameter) is homogeneously distributed in the matrix (
According to the alloy stoichiometries in Table 3, the Alloy 2, Alloy 5 and Alloy 52 were weighed out from high purity elemental charges. 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 2 plates for each alloy under identical processing conditions with nominal dimensions of 65 mm by 75 mm by 1.8 mm thick. The resultant plates were subjected to HIP cycle with subsequent heat treatment. Corresponding HIP cycle and heat treatment for each alloys are listed in Table 16. 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.
Tensile specimens were cut out from each plate that were tested in tension on an Instron mechanical testing frame (Model 3369). The tensile stress-strain curves for Alloy 2, Alloy 5 and Alloy 52 after different annealing are shown in
Using modified tensile specimens with extended grip area, elastic modulus was measured for selected alloy listed in Table 17 in different conditions. Elastic modulus in Table 17 is reported as an average value of 5 separate measurements. As it can be seen, modulus values vary in a range from 192 to 201 GPa depending on alloys chemistry and thermal mechanical treatment.
Using high purity elements, 35 g alloy feedstocks of the Alloy 51 representing Class 2 steel was weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 of 1100° 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 plates were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time.
Tensile specimens were cut out of the plates from the selected alloy which were annealed at 700° C. for 1 hour with air cooling. Annealed specimens were tested in tension on an Instron mechanical testing frame (Model 3369) with recording strain hardening coefficient (n) values as a function of straining during testing utilizing Instron's Bluehill control and analysis software. The results are summarized in
Using high purity elements, 35 g alloy feedstocks of the Alloy 6 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 of 1100° 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 plates were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time. Annealing at 700° C. for 1 hour with slow cooling was applied to plates after HIP cycle. 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.
Tensile specimens were cut out of the plates from the selected alloy which were annealed at 700° C. for 1 hour with slow cooling. Annealed specimens were tested in tension on an Instron mechanical testing frame (Model 3369) with recording strain hardening coefficient (n) values during testing utilizing Instron's Bluehill control and analysis software. A dependence of strain hardening coefficient on tensile strain (elongation) is illustrated in
Using high purity elements, 35 g alloy feedstocks of the Alloy 51 representing Class 2 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 51 was subjected to HIP cycle at 1100° 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 plate was 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.
Tensile specimens were cut out of the plates which were annealed at 850° C. for 1 hour with air cooling. 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 3%. The resultant stress-strain curves are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 6 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 from the alloy were subjected to HIP cycle at 1100° 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 plates 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.
Tensile specimens were cut out of the plates from the selected alloy which were annealed at 700° C. for 1 hour with slow cooling. 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 1%. The resultant stress-strain curves are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 51 representing Class 2 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 51 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 plate was heated at 10° C./min until the target temperature of 1100° C. was reached and was exposed to an isostatic pressure of 30 ksi for 1 hour.
Tensile specimens were cut out of the plates which were annealed at 850° C. for 1 hour with air cooling. 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. Tensile specimen was pre-strained to 10% with subsequent unloading and then tested again up to failure. The resultant stress-strain curves are shown in
SEM images of microstructure in the specimen before and after pre-straining to 10% are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 6 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 6 was subjected to a HIP cycle C (at 1100° 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 plates 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. Tensile specimens were cut from the treated plate.
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. One specimen of the Alloy 6 after HIP cycle at 1100° C. for 1 hour was tested to failure. Another specimen from the same plate was pre-strained to 3%, unloaded and then tested again to failure. The resultant stress-strain curves are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 51 representing Class 2 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 51 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 plates 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. 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. One specimen of the Alloy 51 after HIP cycle at 1100° C. for 1 hour was tested to failure. Another specimen from the same plate was pre-strained to 10%, unloaded, annealed at 1100° C. for 1 hour and then tested again to failure. The resultant stress-strain curves are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 6 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 6 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 plates 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. Tensile specimens were cut from the plate. 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. One specimen of the Alloy 6 after HIP cycle at 1100° C. for 1 hour was tested to failure. Another specimen from the same plate was pre-strained to 3%, unloaded, annealed at 1100° C. for 1 hour and then tested again to failure. The resultant stress-strain curves are shown in
SEM images of microstructure of the gage section of the tensile specimens from Alloy 6 plate (HIPed at 1100° C. for 1 hour and heat treated at 700° C. for 1 hour with slow furnace cooling) tested to failure after pre-straining to 3% and annealing at 1100° C. for 1 hour are shown in
Using high purity elements, 35 g alloy feedstocks of the Alloy 51 representing Class 2 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 51 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 plate was heated at 10° C./min until the target temperature of 1100° C. was reached and was exposed to an isostatic pressure of 30 ksi for 1 hour.
Tensile specimens were cut out of the plates which were annealed at 850° C. for 1 hour with air cooling. 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. The specimen was pre-strained to 10% with subsequent annealing at 1100° C. for 1 hour. Then it was deformed to 10% again twice with subsequent unloading and annealed at 1100° C. for 1 hour. The tensile curves for 3 rounds of pre-straining and testing to failure are shown in
For more detailed structural analysis, TEM specimens were prepared from the grip and from the gage sections of the specimen after cycling deformation. TEM specimens were made first by mechanical grinding/polishing, and then electrochemical polishing. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV. TEM images are presented in
Using high purity elements, 35 g alloy feedstocks of the Alloy 6 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into ingots using high purity argon as a shielding gas. The ingots were flipped several times and re-melted to ensure homogeneity. The resulting ingots 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 plate from the Alloy 6 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 plates 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. Tensile specimen was cut from the plate and heat treated at 700° C. for 1 hour with slow furnace cooling. The tensile specimen was pre-strained to 3% with subsequent annealing at 1100° C. for 1 hour. Then it was deformed to 3% again twice with subsequent unloading and annealed at 1100° C. for 1 hour. The tensile curves for 3 rounds of pre-straining and testing to failure are shown in
The study was performed to evaluate formability of the alloys described in this application at elevated temperatures. In a case of plate production by Twin Roll Casting or Thin Slab Casting, utilized alloys should have good formability to be processed by hot rolling as a step at production process. Moreover, hot forming ability is a critical feature of the high strength alloys in terms of their usage for part production with different configuration by such methods as hot pressing, hot stamping, etc.
Using high purity elements, 35 g alloy feedstocks of the Alloy 20 and Alloy 22 representing Class 3 steel were weighed out according to the atomic ratios provided in Table 3. 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. The resulting ingots 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 plates with thickness of 1.8 mm.
Each resultant plate from the selected alloys was subjected to a HIP cycle specified in Table 18 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 plates were heated at 10° C./min until the target temperature specified for each plate in Table 18 was reached and were exposed to an isostatic pressure of 30 ksi for 1 hour. Heat treatment specified in Table 18 for each plate was applied after HIP cycle. Tensile specimens with a gage length of 12 mm and a width of 3 mm were cut from the treated plates.
The tensile measurements were done at strain rate of 0.001 s−1 at temperatures specified in Table 18. In Table 19, a summary of the tensile test results including total tensile elongation (strain), yield stress, ultimate tensile strength, and location of the failure are shown for the treated plates from Alloy 20 and Alloy 22. Room temperature tensile property ranges for the same alloy after the same treatments are listed for comparison. As can be seen, high strength alloys with ultimate strength up to 1650 MPa at room temperature show high ductility at elevated temperatures (up to 88.5%) demonstrating high hot forming ability. High temperature ductility of the alloys strongly depends on alloy chemistry, thermal mechanical treatment parameters and testing temperature. An example of tested specimen is shown in
Microstructure of the gage of selected specimens from Alloy 20 and Alloy 22 representing Class 3 steel and tested in tension at elevated temperatures as described in Case Example #19, were examined both by SEM and TEM. Samples that were cut out from the gage of the tested specimens 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 manufactured by Carl Zeiss SMT Inc. Example SEM backscattered electron micrographs taken from the gages of tested specimens are shown in
Much less cavitation was observed in the Alloy 22 gage specimens (
TEM was used to characterize the detailed microstructure after the high temperature deformation in the specimens from both alloys. TEM specimens were prepared from the gage of the specimens after high temperature tests until failure. The samples were cut from the tensile gage, 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. These specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope operated at 200 kV.
Energy dispersive spectrometry (EDS) was utilized to characterize the composition in the nano-precipitates. To compare the difference, both the nano-precipitates and matrix are probed by EDS. In
In Alloy 22 specimens, no nano-precipitates were found as compared to that in Alloy 20 specimens. Alloy 22 does not contain copper. However, grain refinement through phase transformation occurred in Alloy 22 specimens tested at both 700° C. and 850° C. The extent of grain refinement is much larger at 700° C. than at 850° C.
The chemistries listed in Table 20 have been used for material processing through plate casting in a Pressure Vacuum Caster (PVC). Using ferroadditives and other readily commercially available constituents, 35 g commercial purity (CP) feedstocks were weighed out according to the atomic ratio provided in Table 20. 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. The resulting ingots were then placed in a PVC chamber, melted using RF induction and then ejected into a copper die designed for casting 3 by 4 inches plates with thickness of 1.8 mm mimicking alloy solidification into plate with similar thickness between rolls at Stage 1 of Twin Roll Casting process.
Thermal analysis was done on the as-solidified cast plate samples on a NETZSCH DSC 404F3 PEGASUS V5 system. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) were performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultra-high purity argon. DTA results are shown in Table 21 indicating the melting behavior for the alloys. As can be seen from the tabulated results in Table 21, the melting occurs in 1 or 2 stages with initial melting observed from ˜1114° C. depending on alloy chemistry. Final melting temperature is up to ˜1380° C. Variations in melting behavior may also reflect complex phase formation at chill surface processing of the alloys depending on their chemistry.
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 22 and was found to vary from 7.63 g/cm3 to 7.66 g/cm3. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3.
Each plate 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 plates were heated at 10° C./min until the target temperature was reached and were exposed to gas pressure for specified time which was held for 1 hour for these studies. HIP cycle parameters are listed in Table 23. The key aspect of the HIP cycle was to remove macrodefects such as pores and small inclusions by mimicking hot rolling at Stage 2 of Twin Roll Casting process or at Stage 1 or Stage 2 of Thin Slab Casting process.
The tensile specimens were cut from the plates after HIP cycle 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 24, a summary of the tensile test results including total tensile elongation (strain), yield stress, and ultimate tensile strength are shown for the cast plates after HIP cycle. Additional column is added that specifies the alloy mechanical response in correspondence with the class of behavior (
After HIP cycle, the plate material was heat treated in a box furnace at parameters specified in Table 25. The key 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. 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.
The tensile specimens were cut from the plates 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 26, a summary of the tensile test results including total tensile elongation (strain), yield stress, and ultimate tensile strength are shown for the cast plates after HIP cycle and heat treatment. Additional column is added that specifies the alloy mechanical response in correspondence with the class of behavior (
High strength/high ductility property combination in the alloys with Class 2 behavior related to the formation of NanoModal Structure (Structure #2,
Using high purity elements, feedstocks with different mass of the Alloy 6 were weighed out according to the atomic ratios provided in Table 3. 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 plate with thickness of 1 inch. Note that the plate that was cast was much thicker than the previous 1.8 mm plates and illustrate the potential for the chemistries in Table 3 to be processed by the Thin Slab Casting process.
The thick plate was cut in half. One part was held in as-cast state. The second part was subjected to HIP cycle at 1000° C. 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 plate was heated at 10° C./min until the target temperature of 1000° C. was reached and was exposed to an isostatic pressure of 30 ksi for 1 hour. Thin plates with thickness of 2 mm were cut from the thick plate in as-cast and HIPed conditions. Three thin plates were cut from the plate after the HIP cycle, which were heat treated at different parameters specified in Table 27. Tensile specimens then were cut from these thin plates in as-cast and HIPed/heat treated conditions. Examples of the partial plate (A), a thin plate from the plate (B) and tensile specimens (C) are shown in
The tensile specimens were cut from the plate 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 27, a summary of the tensile test results including total tensile elongation (strain), yield stress and, ultimate tensile strength is shown for 1 inch thick plate in as-cast state and after HIP cycle with subsequent heat treatments. As can be seen, the tensile strength values vary from 729 to 1175 MPa. The total elongation value varies from 0.49 to 1.05%. Tensile strength and ductility are also illustrated in
The alloys herein in either forms as Class 2 or Class 3 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 2 and/or Class 3 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), pipe casing, 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. The alloys herein may also be used for a compressed gas storage tank and liquefied natural gas canisters.
Class 2 alloys have demonstrated relatively high ductility (up to 25%) at room temperature confirming their cold formability and with further development are expected to reach ductilities up to 40%. Class 3 steels are applicable for various hot forming processes and with further development cold forming applications as well.
This application claims the benefit of U.S. Provisional Application Ser. No. 61/583,261 filed Jan. 5, 2012 and U.S. Provisional Application Ser. No. 61/604,837 filed Feb. 29, 2012.
| Number | Date | Country | |
|---|---|---|---|
| 61583261 | Jan 2012 | US | |
| 61604837 | Feb 2012 | US |
| Number | Date | Country | |
|---|---|---|---|
| Parent | 13862825 | Apr 2013 | US |
| Child | 13863911 | US | |
| Parent | 13556410 | Jul 2012 | US |
| Child | 13862825 | US |
