Patent Grant
8257512
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.
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.
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.
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.
a) illustrates stain hardening in alloy sheets with different mechanisms of structural formation.
b) illustrates tensile properties for the sheets in
a) is a stress-strain curve for Alloy 19 sheet after prestraining to 10%.
b) is a stress-strain curve for Alloy 19 sheet after prestraining to 10% and subsequent annealing at 1150° C. for 1 hour.
a) is a plane view of the plate of Alloy 3 after Erichsen test stopped at maximum load.
b) is a side view of the plate of Alloy 3 after Erichsen test stopped at maximum load.
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).
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. 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.
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
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=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.
Class 2 Steel
As shown in
In
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 (
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
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.
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.
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.
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.
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.
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.
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
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
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.
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.
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.
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 (
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
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
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
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 (
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
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
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
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
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
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
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.
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
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
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
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.
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
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
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
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
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
As it was shown, the Alloy 14 sheet has demonstrated Structure #1 Modal Structure (Step#1) in as-cast state (
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
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
As shown in
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
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.
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
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
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
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
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.
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 (
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
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.
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
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
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
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
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
High strength in pre-strained specimen (
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
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.
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
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.
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 (
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
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.
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.
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.
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.
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.
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
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.
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.
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.
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
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.
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.
| Number | Name | Date | Kind |
|---|---|---|---|
| 4297135 | Giessen et al. | Oct 1981 | A |
| 4576653 | Ray | Mar 1986 | A |
| 6689234 | Branagan | Feb 2004 | B2 |
| 6767419 | Branagan | Jul 2004 | B1 |
| 7323071 | Branagan | Jan 2008 | B1 |
| 8133333 | Branagan et al. | Mar 2012 | B2 |
| Number | Date | Country | |
|---|---|---|---|
| 61586951 | Jan 2012 | US | |
| 61488558 | May 2011 | US |
