JP2017509802A - A new class of high-performance, high-strength steel that can be warm formed - Google Patents

Abstract

Fe of 48.0 atomic% to 81.0 atomic%, B of 2.0 atomic% to 8.0 atomic%, Si of 4.0 atomic% to 14.0 atomic%, and Cu, Mn or Ni At least one of them, Cu is present at 0.1 atomic percent to 6.0 atomic percent, Mn is present at 0.1 atomic percent to 21.0 atomic percent, and Ni is 0.1 atomic percent To 16.0 atomic percent of metal alloys are disclosed. The alloy can be heated at temperatures from 200 ° C. to 850 ° C. for times up to 1 hour, and there is no eutectoid transformation on cooling. The alloy can then be formed into a selected shape.

Description

This application claims priority to US Provisional Application No. 61 / 768,131, filed February 22, 2013.

FIELD OF THE INVENTION The present disclosure is directed to a new type of high performance high strength steel (AHSS) that can be warm formed. This steel can be warm formed due to its inherent structure, allowing it to develop to a relatively high strength without the need for austenitization and quenching.

  Existing hot-formed steel is a variant of martensite grate manufactured under various trade names including USIBOR®, DUXTIBOR® and the like. This class of materials can typically develop strengths at 1200 to 1600 MPa with limited ductility of 5-8%. In the as-manufactured state, these grades of steel are in their annealed soft state and are composed primarily of ferrite and cementite and therefore exhibit low tensile strength. In order to produce high strength parts, the steel must then be heated to its austenitizing temperature (ie A3), which is typically in the range of 850 to 1000 degrees, depending on the chemical nature. After an appropriate holding time to form a single phase solid solution of austenite, the steel is then deformed to produce parts that can be a wide variety of structural and non-structural components. After deformation, the part is held to ensure that the shape is maintained and then quenched in oil or water depending on the thickness of the part being formed and the inherent hardenability of the steel alloy. . Often, small additions of boron, typically up to 0.05 wt%, are used to improve the hardenability of the steel, which means that the process window for martensite formation is widened. With proper quenching, the steel part then forms a strong and fragile martensitic structure. Subsequent heat treatment is typically performed to produce tempered martensite that provides improved ductility by sacrificing some strength level.

  The present disclosure is directed to a steel alloy that can be warm formed (treated at a temperature of 200 ° C. to 850 ° C. for a time of 1.0 second to 1 hour by either direct heating or induction heating). . The elemental composition range (atomic%) is: Fe present from 48.0 to 81.0, B from 2.0 to 8.0, Si from 4.0 to 14.0, and Cu, Mn and Ni At least one austenite stabilizer (an element that stabilizes austenite formation) including one or more of Cu, 0.1 to 6.0 atomic% Cu, and 0.1 to 21.0 atomic% Mn. Present, and Ni is present at 0.1 to 16.0 atomic percent. Optionally, Cr can be included at a level of up to 32.0 atomic%. Any other element such as C, Al, Ti, V, Nb, Mo, Zr, W and Pd may be present up to 10.0 atomic percent. Impurities expected / known to be present include Nb, Ti, S, O, N, P, W, Co, Sn, and may be present at levels up to 10.0 atomic percent. Alloys herein suitable for warm forming include Class 1, Class 2 and Class 3 steels described herein. The steel alloys of the present disclosure having application to centrifugal casting are in a wide range of ductility and strength, depending on the above class of steel, due to the new effective structure type promoted by a new effective mechanism. Provide a unique combination of properties.

It is a binary phase diagram regarding the iron rich region of the iron-carbon binary system. It is a binary Fe-C phase diagram showing the difference between a new grade of warm-formed steel (top blowout) and conventional steel (bottom blowout). FIG. 3 is a model phase diagram showing an expected phase equilibrium of a new warm-formed steel type. The structure and mechanism for the formation of class 1 steel herein is shown. 2 shows a typical stress-strain curve for a material having a modal structure. The structure and mechanism for the formation of the class 2 steel alloy herein is shown. Figure 2 shows stress-strain curves for the relevant mechanism and class structure shown in class 2 alloys. The structure and mechanism for the formation of the class 3 steel alloy herein is shown. Figure 2 shows stress-strain curves for the relevant mechanism and class structure shown in class 3 alloys. It is a photograph of the plate in an as-cast state. A nanosteel size R & D specimen shape modified to increase the grip to 9.5 mm to accommodate a 1/8 "grip pinhole. Temperature dependence of tensile elongation and yield stress in alloy 213. Photos of 36 class 3 alloy specimens after heat treatment and HIP cycle before and after deformation to 57.5%. Tensile strength, yield stress and tensile elongation as a function of test temperature in commercial sheets from alloy 82.

A new class of warm-formed steels A new class of warm-formed steels does not need to be austenitized due to very different metallurgy, enabling a metallurgical transformation (ie austenite to martensite Not) In FIG. 1, the iron rich binary component of the binary Fe—C phase diagram is shown. This figure is used to illustrate the basic phase equilibria in ˜30,000 known worldwide equivalent iron and steel alloys. In FIG. 2, the Fe—C binary phase diagram is utilized to show the difference between a new class of warm-formed steel and conventional steel. With the exception of austenitic stainless steels and TWIP (twinning induced plasticity) steels, almost all conventional steels have evolved by focusing primarily on structural development and heat treatment based on eutectoid transformation. While the heat treatment temperature, time and strategy can vary widely, in general, the first step is to heat the steel to the single phase austenite region. Since the hardenability of steel is sensitive to the average particle size of the material, the heating rate to the target temperature and the time at the temperature are important. Depending on how the steel is cooled or quenched from the austenitizing temperature, a wide range of manufactured characteristic structures including pearlite, upper and lower bainite, spheroidite, and martensite. Will bring. In addition, complex or biphasic microstructures can be produced with different fractions of all these characteristic microstructures, together with ferrite, residual austenite and cementite phases.

  As shown in FIG. 2, the new class of warm formed steel is essentially different because the focus on phase and structure development is on the peritectic region rather than the eutectoid region. It should be noted that the peritectic invariant reaction includes a liquid comprising delta ferrite + native transformation liquid that produces austenite. This is very different from the solid state eutectoid transformation containing austenite that produces ferrite and cementite.

  To further illustrate these differences, a model phase diagram for warm-formed alloys is provided in FIG. The x-axis (labeled as atomic% alloying) is a reference to an alloy comprising Fe, B and Si and at least one of Cu, Mn or Ni as described above. And the temperature on the y-axis will vary depending on the alloy selected. As can be seen, the eutectoid transformation, which is critical to existing steels, lacks the complex multicomponent phase diagram for the steels herein. The transition involves an early austenite-to-ferrite transformation associated with the gamma / austenite stability loop and initial solidification via peritectic transformation.

  New types of steel produced herein may include any of the class 1, class 2 or class 3 steel alloys described herein above that are warm formed, but preferably Includes warm forming of Class 2 or Class 3 steel alloys. These Class 1, Class 2 and Class 3 steel structures are stable to high temperatures and hot formed at conventional temperatures known for hot forming processes with typical hot forming ductility of 30 to 120%. Can be done. However, Class 1, Class 2 and Class 3 steels herein exhibit relatively high strength and ductility at room temperature and maintain that high ductility at warm temperatures (ie, 200 to 850 °). Therefore, it is applicable regarding cold deformation through various methods including cold rolling, stamping, roll forming, hydroforming and the like. Furthermore, class 1, class 2 and class 3 steels can now be processed by a warm forming process. In warm forming, the steel described above is now in a temperature range that is less than hot forming for a period of time of 1 second to 1 hour via direct heating (eg furnace heating) and / or induction heating, typically Is heated from 200 to 850 ° C. This temperature range allows manufacturing for a number of major factors that will be described later. In short, warm forming can now reduce costs while creating new functionality through minimizing or avoiding the springback problems found in cold formed steel.

Possible advantages of warm-formed steel / new functionality Zinc coating Steel provides an anode sacrificial coating to protect the surface of the steel from corrosion, commonly through a process called galvanization. Protected from corrosion. There are various ways to apply zinc or zinc alloys to the surface, including conventional galvanizing, hot dip galvanizing, galvannealing and the like. All of these processes share similar characteristics with zinc bound to different degrees to the steel surface. This is a problem with hot forming. This is because zinc exhibits a low melting point of 419 ° C. As such, during hot forming of conventional martensite / press formable steel, the zinc coating melts and evaporates, leaving the resulting steel part vulnerable to corrosion attack. Efforts have been made to shorten the hot forming cycle time to produce a thicker initial layer of zinc and / or to limit high temperature exposure, but the results are ineffective. There is a costly post part forming coating step to restore the anode surface. Through warm forming at temperatures below the melting point of zinc (ie, ~ 200-419 ° C), the problem of zinc loss can be minimized or completely avoided. Therefore, the new nanomodal steel processed via warm forming creates new functionality through the ability to precoat by conventional galvanizing processes and then this in the finished warm deformed parts Maintain protective coating.

Cycle time Traditional hot forming lines have conveyor type continuous ovens that allow hot formed parts to be fed in a continuous fashion that reaches their target austenitizing temperature prior to hot deformation. Use. The length of these continuous gas fired ovens can be 50 meters or more, and if any problem occurs during the hot forming operation, all parts moving through a long furnace are generally Discarded. This is because during subsequent reheating, their metallurgical structures will be adversely and irreversibly affected. By heating to a lower temperature for warm forming, the length of this continuous oven used will need to be very short, so less infrastructure, lower amounts Require waste parts, and especially lower energy costs. This ultimately results in lower cost parts, thus enabling the technology for a wide range of applications.

Oxidation / Post Treatment A cost factor that limits hot forming is scale / oxide removal that forms during high temperature exposure and then needs to be removed by existing shot / grit blasting. Oxidation occurs due to the high temperature exposure essential to austenitize existing materials. Furthermore, the process does not give itself an inert gas atmosphere. This is because after hot forming, the part must be quenched in a liquid medium to form martensite, which produces additional oxidation. With a new class of warm-formed steel, the temperature of deformation will be much lower and will limit / prevent typical oxidation with high temperature exposure. In addition, warm-formed steels do not need to be quenched and they show an unreactive response to the cooling rate in the solid state so that warm-formed parts are processed while maintained in an inert atmosphere. It may be possible to prevent or minimize oxidation. And since scale formation is avoided, this is expected to result in parts that do not need to go through expensive grit / shot blasting.

Cooling / Water Quenching Existing hot formed steels need to be quenched from their high austenitizing temperature to form a martensitic structure that provides high strength. During quenching into oil, water, brine brine, etc., parts can be distorted and / or cracked, resulting in higher waste rates. In addition, the formation of martensite texture is highly dependent on the cooling rate, so that some areas of insufficient cooling can occur, for example when a vapor barrier is generated from a liquid medium. This results in lower strength levels in certain areas, often resulting in higher gauge thickness and higher weight than is essential, while being considered in component design to overcome local strength variations Generated strength debt. A new class of nanomodal warm-formed steel need not be water quenched and need not be heated to the high temperatures found in conventional austenitization. Therefore, strict dimensional control is possible due to the lack of quenching deformation. This results in a lower waste rate and reduced costs, enabling the technology.

Preforming / final finish Due to the fact that existing martensitic steels need to be austenitized at high temperatures, hot deformed and then quenched in liquid media, the resulting parts are originally It is distorted from the blank dimension. Due to the presence of distortion, the final details in the part (ie final trimming, hole incorporation, etc.), especially during quenching, can be pre-shaped in the starting blank. Therefore, a mechanical re-hit in a post stamping operation, or expensive laser trimming is required, and the final hole insertion for trimming to the final part dimensions. It requires expensive molds that require regular maintenance to handle the extremely strong materials obtained from the hot forming required as a post-finishing process. There is a fairly low temperature range that results in significantly less thermal expansion through warm forming, combined with the lack of need for quenching, which makes warm-formed steel a design and previously unknown. It means presenting process capability. As such, the starting blank can be properly trimmed prior to warm forming and fully or partially pre-formed by holes, creating new functionality and creating a final finish inherent in existing hot forming processes. Eliminating costly laser trimming processes.

New Class Steel Alloys The non-stainless steel alloys herein are present as Class 1, Steel, Class 2 or Class 3 steels, which are preferably crystalline (non-glass) with a distinct grain size morphology. It is such that it allows the formation of what is described in the specification. The ability of alloys to form Class 1, Class 2 or Class 3 steels herein is described in detail herein. However, it is useful to first consider the description of the general features of class 1, class 2 and class 3 steels provided below.

Class 1 Steel The formation of Class 1 steel herein is shown in FIG. As shown there, a modal structure is first formed, which is the result of starting with a liquid melt of the alloy and solidifying by cooling, with nucleation and a specific particle size Provide specific phase growth. Thus, references herein to modal can be understood as structures having at least two particle size distributions. The particle size herein can be understood as the size of a single crystal of a specific specific phase, preferably distinguishable by methods such as scanning electron microscopy or transmission electron microscopy. Thus, Class 1 steel structure # 1 can preferably be achieved by processing either through laboratory scale procedures and / or industrial scale methods such as powder atomization or alloy casting.

Thus, the modal structure of class 1 steel will initially exhibit the following particle sizes when cooled from the melt: (1) 500 nm to 20,000 nm containing austenite and / or ferrite Matrix particle size; (2) Boride particle size of 25 nm to 500 nm (that is, non-metallic particles such as M ₂ B, M is a metal and is covalently bonded to B). The boride grains may also preferably be a “pinning” type phase, which is a reference to the feature that the matrix grains are effectively stabilized by a pinning phase that resists coarsening at high temperatures. Metal boride grains have been identified as exhibiting M ₂ B stoichiometry, but other stoichiometry is possible, and M ₃ B, MB (M ₁ B ₁ ), M ₂₃ B ₆ and M ₇ B ₃ it should be noted that to provide a pinning including.

  Class 1 steel modal structures may undergo thermomechanical deformation and / or heat treatment, resulting in some changes in properties, but the modal structure can be maintained.

  When the above class 1 steel is exposed to mechanical stress, the observed stress versus strain diagram is shown in FIG. Thus, it is observed that the modal structure leads to a second type of structure for class 1 steel, which is what is identified as dynamic nanophase precipitation, which is a modal nanophase structure. Thus, such dynamic nanophase precipitation is caused when the alloy experiences yield under stress, and the yield strength of class 1 steels undergoing dynamic nanophase precipitation can preferably occur between 400 MPa and 1300 MPa. Have been found. Thus, it can be seen that dynamic nanophase precipitation occurs due to the application of mechanical stress above such indicated yield strength. Dynamic nanophase precipitation itself can be understood as the formation of a further identifiable phase in class 1 steel, referred to as a precipitated phase with an associated particle size. In other words, the result of such dynamic nanophase precipitation is that the discriminating matrix particle size from 500 nm to 20,000 nm, from 25 nm, together with the formation of precipitated grains containing hexagonal phase and grains from 1.0 nm to 200 nm. The formation of an alloy that still exhibits a boride pinning particle size of 500 nm. Thus, as described above, the grain size does not coarsen when the alloy is pressured, leading to the development of precipitated grains as described.

Reference to the hexagonal phase is ditrigona with a dihexagonal pyramid class hexagonal phase with P6 ₃ mc space group (# 186) and / or a hexagonal P6bar2C space group (# 190). It can be understood as a dipyragonal dipyramidal class. In addition, the mechanical properties of such a second type of structure of class 1 steel are such that the tensile strength is observed to fall in the range of 700 MPa to 1400 MPa with 10-50% elongation. Is. Furthermore, the second type of structure of class 1 steel is such that it exhibits a strain hardening coefficient of 0.1 to 0.4 that is substantially flat after undergoing the indicated yield. The strain hardening factor is a reference to the n value in the formula σ = Kε ⁿ , where σ represents the applied stress on the material, ε is the strain, and K is the strength factor. The value of the strain hardening index n is between 0 and 1. A value of 0 means that the alloy is a completely plastic solid (ie, the material undergoes irreversible changes to the applied force), while a value of 1 is a 100% elastic solid (ie, The material undergoes a reversible change with respect to the applied force).

  Table 1A below provides a comparison and performance summary for the class 1 steels herein.

Class 2 Steel The formation of Class 2 steel herein is shown in FIG. Class 2 steels may also be formed herein from the identified alloys and are described herein as static nanophase refinement and dynamic nanophase strengthening after starting with structure type # 1, modal structure. It contains two new structure types followed by two new mechanisms identified in. New structural types for class 2 steels are described herein as nanomodal structures and high strength nanomodal structures. Accordingly, the class 2 steels herein can be characterized as follows: Structure # 1-modal structure (step # 1), mechanism # 1-static nanophase refinement (step # 2), structure # 2-Nanomodal structure (step # 3), mechanism # 2-dynamic nanophase strengthening (step # 4), and structure # 3-high strength nanomodal structure (step # 5).

  As shown there, structure # 1 is formed first, and the modal structure is the result of starting with the liquid melt of the alloy and solidifying by cooling, which results in nucleation and specific particle size. Provide specific phase growth to have. The particle size herein can be understood again as the size of an intrinsic specific phase single crystal that is preferably distinguishable by methods such as scanning electron microscopy or transmission electron microscopy. Thus, class # 2 steel structure # 1 may preferably be achieved by any treatment via laboratory scale procedures and / or via industrial scale methods such as powder atomization or alloy casting.

Thus, the modal structure of Class 2 steel will initially exhibit the following particle sizes when cooled from the melt: (1) 500 nm to 20,000 nm matrix particle size containing austenite and / or ferrite (2) a boride particle size of 25 nm to 500 nm (ie, non-metallic particles such as M ₂ B, where M is a metal and is covalently bonded to B). The boride grains may also preferably be a “pinning” type phase referring to the feature that the matrix grains will be effectively stabilized by a pinning phase that resists coarsening at high temperatures. Metal boride grains have been identified as exhibiting M ₂ B stoichiometry, but other stoichiometry is possible, and M ₃ B, MB (M ₁ B ₁ ), M ₂₃ B ₆ and M ₇ can provide a pinning including B _3, it should be noted that not affected by the mechanism # 1 or # 2 above. Reference to particle size will again be understood as the size of a single crystal of a specific specific phase, preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Furthermore, class 2 steel structure # 1 herein includes austenite and / or ferrite in combination with such boride phases.

  In FIG. 7, a stress strain curve is shown and represents a non-stainless steel alloy herein subject to the deformation behavior of class 2 steel. The modal structure is preferably generated first (structure # 1), after which the modal structure is now uniquely refined through mechanism # 1, which is a static nanophase refinement mechanism that leads to structure # 2. Can be done. Static nanophase refinement is a matrix grain that typically falls in the range of 100 nm to 2000 nm as the matrix particle size of structure 1 first falls in the range of 500 nm to 20,000 nm, but decreases in size. Reference is made to the feature of providing a structure 2 having a diameter. It should be noted that while the boride pinning phase can vary significantly in size in some alloys, it is designed to resist matrix grain enlargement during heat treatment. Due to the presence of these boride pinning sites, grain boundary movement leading to coarsening would be expected to be delayed by a process called Zener pinning or Zener drag. Therefore, matrix grain growth may be energetically favorable due to a decrease in total interfacial area, while the presence of boride pinning phases is coarsening due to the high interfacial energy of these phases. This driving force will be countered.

  A feature of the static nanophase refinement mechanism in class 2 steels, the micron-scale austenite phase (gamma-Fe), described as falling within the 500 nm to 20,000 nm range, is a new phase (eg, ferrite or Transform partially or completely to alpha-Fe). The volume fraction of ferrite (alpha-iron) initially present in the modal structure of class 2 steel (Structure 1) 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 20 to 80%. The static transformation preferably includes an inherent refinement mechanism since it occurs during high temperature heat treatment. This is because grain coarsening rather than grain refinement is a conventional material response at high temperatures.

  Thus, grain coarsening does not occur with the class 2 steel alloys herein during the static nanophase refinement mechanism. Structure # 2 can uniquely transform to structure # 3 during dynamic nanophase strengthening, resulting in the formation of structure # 3 and in the range of 800 to 1800 MPa with a total elongation of 5 to 40% The tensile strength value in is shown.

  Depending on the chemistry of the alloy, nanoscale precipitates can form during static nanophase refinement and subsequent thermal processes in some of the non-stainless high strength steels. The nanoprecipitates range from 1 nm to 200 nm and the majority of these phases (> 50%) are 10-20 nm in size and formed in structure # 1 with respect to slowing the coarsening of the matrix grains Much smaller than the boride pinning phase. Also, during static nanophase refinement, the boride particle size grows larger to sizes in the range of 200 to 2500 nm.

  More specifically, in the case of the alloys herein that provide class 2 steel, when such alloys exceed their yield point, plastic deformation at a constant stress occurs, after which structure # 3 The dynamic phase transformation that leads to generation continues. More specifically, after sufficient stress is induced, the slope of the stress versus strain curve changes to increase the inflection point (FIG. 7), the strength increases with strain, and mechanism # 2 (dynamic Activation of nanophase enhancement).

  With further strain during dynamic nanophase strengthening, the strength continues to increase, but the strain hardening coefficient gradually decreases until almost at break. Some strain softening occurs, but only near the break point that can be attributed to a decrease in local cross-sectional area at necking. It should be noted that the strengthening transformation that occurs due to material strain under stress generally defines mechanism # 2 as a dynamic process leading to structure # 3. By dynamic is meant that the process can occur by the application of stresses beyond the yield point of the material. Tensile properties that can be achieved for alloys that achieve structure 3 include tensile strength values in the range of 800 to 1800 MPa and total elongation of 5 to 40%. The level of tensile properties achieved also depends on the amount of transformation that occurs when the strain increases corresponding to the characteristic stress strain curve for class 2 steel.

  Therefore, depending on the level of transformation, an adjustable yield strength can also now develop in the class 2 steels of this specification, depending on the level of deformation, and in structure # 3, the yield strength is finally from 400 MPa. It can vary up to 1700 MPa. That is, conventional steels outside the scope of the alloys herein show only a relatively low level of strain hardening, so their yield strength can be in a small range (e.g., 100 to 200 MPa) depending on prior deformation history. ) Only. For class 2 steels herein, yield strength can vary over a wide range (eg, 400 to 1700 MPa) when applied to the transformation of structure # 2 to structure # 3, with adjustable deformation varying It gives potential to both designers and end users in the application, and makes structure # 3 available for various applications such as collision management in car body structures.

With this dynamic mechanism, new and / or additional precipitated phases or multiple precipitated phases are observed, showing discernable particle sizes from 1 nm to 200 nm. In addition, there is further distinction in the precipitated phase, dihexagonal pyramidal class hexagonal phase with P6 ₃ mc space group (# 186), ditrigonal dipyramidal class with hexagonal P6bar2C space group (# 190), And / or M ₃ Si cubic phase with Fm3m space group (# 225). Thus, dynamic transformation can occur partially or completely, resulting in the formation of a microstructure with a novel nanoscale / near nanoscale phase that provides a relatively high strength in the material. That is, structure # 3 has matrix grains generally sized between 100 nm and 2000 nm pinned by a boride phase that is in the range of 200 to 2500 nm, and has a precipitated phase that is in the range of 1 nm to 200 nm. It can be understood as a microstructure. The initial formation of the precipitated phase referred to above having a particle size of 1 nm to 200 nm starts with static nanophase refinement and continues during dynamic nanophase strengthening, leading to the formation of structure 3. The volume fraction of the precipitated phase having a particle size of 1 nm to 200 nm in structure 2 increases in structure 3 and is aided by the identified strengthening mechanism. It should also be noted that in structure 3, the level of gamma-iron is arbitrary and can be eliminated depending on the chemistry and austenite stability of the particular alloy.

  It should be noted that dynamic recrystallization is a known process but is different from mechanism # 2. This is because it involves the formation of large grains from small grains, which is a coarsening mechanism rather than a refinement mechanism. In addition, when a new undeformed grain is replaced by a deformed grain, no phase change occurs in contrast to the mechanism given here, which is also the strengthening mechanism here. In contrast, there is a corresponding decrease in intensity. Although metastable austenite in steels is known to transform to martensite under mechanical stress, preferably the evidence for martensite or body-centered tetragonal iron phase is the new described in this application. It should also be noted that it is not found in steel alloys.

  Table 1B below provides a comparison of the structural and performance characteristics of the class 2 steels herein.

Class 3 Steel Class 3 steel is associated with the formation of high strength lamellar nanomodal structures via a multi-step process as will be described herein.

  To achieve a tensile response with sufficient ductility and high strength in non-stainless carbon-free steel alloys, a preferably seven-step process is now disclosed and shown in FIG. Structural development begins with structure # 1-modal structure (step # 1). However, mechanism # 1 in class 3 steel is now related to lath phase generation (step # 2) which leads to structure # 2-modal lath phase structure (step # 3), and mechanism # 2-lamellar nanophase generation (step # 2) Via 4), the structure is transformed into a structure # 3-lamellar nanomodal structure (step # 5). The deformation of structure # 3 results in the activation of mechanism # 3-dynamic nanophase enhancement (step # 6) and leads to the formation of structure # 4-high strength lamellar nanomodal structure (step # 7). Reference is also made to Table 1C below.

Structure # 1, which includes the formation of modal structures (ie bi-, tri- and higher order), can be performed through laboratory scales as shown and / or chill surface treatments such as twin roll casting or thin slab casting. It can be achieved in alloys with the referenced chemistry in this application by processing via industrial scale methods including. Thus, the modal structure of class 3 steel initially exhibits the following particle sizes when cooled from the melt: (1) ferrite or alpha-Fe (essential) and optionally austenite or gamma A matrix particle size of 500 nm to 20,000 nm containing Fe; and (2) a boride particle size of 100 nm to 2500 nm (ie non-metallic particles such as M ₂ B, where M is a metal and B (3) yield strength from 350 to 1000 MPa; (4) tensile strength from 400 to 1200 MPa and total elongation from 0 to 3.0%. It will also show a dendritic growth morphology of the matrix grain. Also, the boride grains may preferably be a “pinning” type phase that refers to the characteristic that the matrix grains will be effectively stabilized by a pinning phase that resists coarsening at high temperatures. Although the metal boride grains have been identified as exhibiting M ₂ B stoichiometry, other stoichiometry is possible, such as M ₃ B, MB (M ₁ B ₁ ), M ₂₃ B ₆ and M It should be noted that pinning involving ₇ B ₃ may be provided and is not affected by mechanism # 1, # 2 or # 3 described above. Reference to particle size will again be understood as the size of a single crystal of a specific specific phase, preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Accordingly, class 3 steel structure # 1 herein includes ferrites in conjunction with such boride phases.

  Structure # 2 includes the formation of a modal lath phase structure with uniformly distributed precipitates from a modal structure (structure 1) with a dendritic morphology through mechanism # 1. The lath phase structure can be generally understood as a structure composed of plate-shaped crystal grains. A reference to “dendritic form” can be understood as a tree, and a reference to “plate-shaped” can be understood as a sheet. The lath structure formation preferably occurs at a high temperature (eg, at a temperature of 700 ° C. to 1200 ° C.) via grain formation such as a plate, and (1) a lath that is typically 100 to 10,000 nm. Structural particle size; (2) boride particle size from 100 nm to 2500 nm; (3) yield strength from 350 MPa to 1400 MPa; (4) tensile strength from 350 MPa to 1600 MPa; (5) 0-12% elongation. . Structure # 2 also contains alpha-Fe, with gamma-Fe remaining as needed.

A second phase of boride precipitate, typically having a size of 100 to 1000 nm, can be found dispersed in the lath matrix as isolated particles. The second phase of the boride precipitate can be understood as non-metallic grains of different stoichiometry (M ₂ B, M ₃ B, MB (M ₁ B ₁ ), M ₂₃ B ₆ and M ₇ B ₃ ). Here, M is a metal and is covalently bonded to boron. These boride precipitates are distinguished from boride grains in structure # 1 having a size that changes little or no change.

Structure # 3 (lamellar nanomodal structure) includes the formation of lamellar morphology as a result of static transformation of ferrite into one or several phases via mechanism # 2 identified as lamellar nanophase formation. Static transformation is the decomposition of the parent phase into a new phase or several new phases due to the partitioning of the alloying elements by diffusion during the high temperature heat treatment, preferably in the temperature range of 700 ° C. to 1200 ° C. Can occur. A lamella (or layered) structure is composed of alternating layers of two phases, whereby individual lamellae are present in colonies connected in three dimensions. For class 3 alloys, lamellar nanomodal structures are: (1) 100 nm to 1000 nm wide lamellae with a thickness in the range of 100 nm to 10,000 nm and with a length of 0.1 to 5 microns; ) 100 nm to 2500 nm boride grains of different stoichiometry (M ₂ B, M ₃ B, MB (M ₁ B ₁ ), M ₂₃ B ₆ and M ₇ B ₃ ), wherein M is a metal, It is covalently bonded to boron; (3) precipitated grains of 1 nm to 100 nm; and (4) yield strength of 350 MPa to 1400 MPa. The lamellar nanomodal structure continues to contain alpha-Fe and gamma-Fe remains arbitrary.

  The lamellar nanomodal structure (structure # 3) is a dynamic nanophase reinforcement (mechanism # 3, mechanical # 3) during plastic deformation (ie, exceeding the yield stress for the material) that exhibits relatively high tensile strength in the 1000 to 2000 MPa range. Transformation to structure # 4 via exposure to adaptive power). In FIG. 9, a stress-strain curve is shown and represents an alloy having structure # 3 herein that undergoes the deformation behavior of class 3 steel compared to that of class 2. As shown in FIG. 9, structure # 3 provides the indicated curve upon application of stress, resulting in structure # 4 of class 3 steel.

  Strengthening during deformation occurs when the material is strained under stress and is related to the phase transformation defining mechanism # 3 as a dynamic process. For alloys for exhibiting high strength at the levels described in this application, the lamellar structure is preferably formed prior to deformation. Inherent in this mechanism is that the micron-scale austenite phase is transformed into a new phase by a decrease in the scale of the microstructural features, typically down to nanoscale morphology. Some proportion of austenite may initially form in some class 3 alloys during casting and then remain present in structure # 1 and structure # 2. During strain when stress is applied, new or additional phases are formed, typically having nanoparticles in the range of 1 to 100 nm.

  In post-deformed structure # 4 (high-strength lamellar nanomodal structure), the ferrite grains comprise alternating layers with nanostructures composed of new phases formed during deformation. Depending on the specific chemistry and the stability of the austenite, some austenite may additionally be present. In structure # 4, a large amount of nanoparticles of different phases are present as a result of dynamic nanophase strengthening, as opposed to the layers in structure # 3, where each layer represents a single or just a few grains. Since nanoscale phase formation occurs during alloy deformation, it represents a stress-induced transformation and is defined as a dynamic process. Nanoscale phase precipitation during deformation causes a large strain hardening of the alloy. Dynamic transformation can occur partially or completely, forming a microstructure with a novel nanoscale / near nanoscale phase identified as a high strength lamellar nanomodal structure (structure # 4) that provides high strength in the material Bring. Thus, structure # 4 can be formed by varying levels of reinforcement depending on the amount of reinforcement achieved by mechanism # 3 and the particular chemistry.

  Table 1C below provides a comparison of the structural and performance characteristics of the class 3 steels herein.

Alloy properties In new alloys, melting occurs in one or more stages, initial melting is from ~ 1000 ° C depending on the alloy chemistry, and final melting temperature is up to ~ 1500 ° C. obtain. Variations in melting behavior reflect complex phase formation in alloy chill surface processing, depending on their chemistry. The density of the alloy varies from 7.2 g / cm ³ to 8.2 g / cm ³ . The mechanical property values in alloys from each class will depend on the chemistry of the alloy and the processing / processing conditions. For class 1 steels, the ultimate tensile strength value can vary from 700 to 1500 MPa with a tensile elongation of 5 to 40%. The yield stress is in the range from 400 to 1300 MPa. For class 2 steels, ultimate tensile strength values can vary from 800 to 1800 MPa with a tensile elongation of 5 to 40%. The yield stress ranges from 400 to 1700 MPa. For class 3 steels, the ultimate tensile strength value can vary from 1000 to 2000 MPa with a tensile elongation of 0.5 to 15%. The yield stress is in the range from 500 to 1800 MPa. Additional classes of steel are expected to have possible yield strength, tensile strength and elongation values outside the limits listed above.

Examples Preferred Alloy Chemistry and Sample Preparation The chemical composition of the investigated alloys is shown in Table 2 which provides the preferred atomic ratios utilized. These chemistries have been investigated by using material processing via sheet casting in a pressure vacuum caster (PVC). Using high purity elements or iron additives and other readily commercially available ingredients, 35 g of the alloy material of the target alloy was weighed according to the atomic ratios provided in Table 2. The raw material was then placed in the copper hearth of the arc melting system. The raw material was arc melted into an ingot using high purity argon as a shielding gas. The ingot was inverted several times and remelted to ensure uniformity. After mixing, the indots were then cast into finger shapes approximately 12 mm wide, 30 mm long and 8 mm thick. The resulting finger is then placed in a PVC chamber, melted using RF induction, and then on a copper mold designed to cast a 3 × 4 inch sheet having a thickness of 1.8 mm. Was released. An example of a cast plate is shown in FIG. Die casting of alloys utilized at relatively high cooling rates that can correlate with metal solidification in different sheet manufacturing processes, including but not limited to sheet solidification on chill surfaces in twin rolls, thin strips and thin slab castings. Related to melting and solidification of

  Therefore, the atomic% of Fe present is 48.0, 48.1, 48.2, 48.3, 48.4, 48.5, 48.6, 48.7, 48.8, 48.9, 49.0, 49.1, 49.2, 49.3, 49.4, 49.5, 49.6, 49.7, 49.8, 49.9, 50.0, 50.1, 50. 2, 50.3, 50.4, 50.5, 50.6, 50.7, 50.8, 50.9, 51.0, 51.1, 51.2, 51.3, 51.4, 51.5, 51.6, 51.7, 51.8, 51.9, 52.0, 52.1, 52.2, 52.3, 52.4, 52.5, 52.6, 52. 7, 52.8, 52.9, 53.0, 53.1, 53.2, 53.3, 53.4, 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, 59.9, 0.0, 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, 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, 73.0, 73.1, 73.2, 73 3, 73.4, 73.5, 73.6, 73.7, 73.8, 73.9, 74.0, 74.1, 74.2, 74.3, 74.4, 74.5, 74.6, 74.7, 74.8, 74.9, 75.0, 75.1, 75.2, 75.3, 75.4, 75.5, 75.6, 75.7, 75. 8, 75.9, 76.0, 76.1, 76.2, 76.3, 76.4, 76.5, 76.6, 76.7, 76.8, 76.9, 77.0, 77.1, 77.2, 77.3, 77.4, 77.5, 77.6, 77.7, 77.8, 77.9, 78.0, 78.1, 78.2, 78. 3, 78.4, 78.5, 78.6, 78.7, 78.8, 78.9, 79, 79.1, 79.2, 79.3, 79.4, 79.5, 79. 6, 79.7, 79.8, 79.9, 80 It may be a 0,80.1,80.2,80.3,80.4,80.5,80.6,80.7,80.8,80.9,81.0.

  Therefore, the atomic% of B is 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. Can be zero.

  Accordingly, the atomic% of Si is 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, 1 .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 obtain.

  Therefore, the atomic% of Cu is 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 .

  Therefore, the atomic ratio of Mn is 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, 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, 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.

  Therefore, the atomic ratio of Ni is 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, 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 It may be a 1,15.2,15.3,15.4,15.5,15.6,15.7,15.8,15.9,16.0.

  Therefore, if present, the atomic ratio of Cr as an arbitrary element is 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, 4.8, 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, 21.1, 21.2, 21.3, 21.4 , 21.5, 21.6, 21.7, 21.8, 21.9, 22.0, 22.1, 22.2, 22.3, 22.4, 22.5, 22.6, 22 7, 22.8, 22.9, 23.0, 23.1, 233.2, 23.3, 23.4, 23.5, 23.6, 23.7, 23.8, 23.9 24.0, 24.1, 24.2, 24.3, 24.4, 24.5, 24.6, 24.7, 24.8, 24.9, 25.0, 25.1, 25 2, 25.3, 25.4, 25.5, 25.6, 25.7, 25.8, 25.9, 26.0, 26.1, 26.2, 26.3, 26.4 , 26.5, 26.6, 26.7, 26.8, 26.9, 27, 27.1, 27.2, 27.3, 27.4, 27.5, 27.6, 27.7. , 27.8, 27.9, 28.0, 28.1 28.2, 28.3, 28.4, 28.5, 28.6, 28.7, 28.8, 28.9, 29.0, 29.1, 29.2, 29.3, 29. 4, 29.5, 29.6, 29.7, 29.8, 29.9, 30.0, 30.1, 30.2, 30.3, 30.4, 30.5, 30.6, 30.7, 30.8, 30.9, 31.0, 31.1, 31.2, 31.3, 31.4, 31.5, 31.6, 31.7, 31.8, 31. It can be 9,32.0.

Case Example # 1: Warm Formability of Class 2 Stainless Alloy A study was conducted to evaluate the warm formability of the alloys described herein at high temperatures. In the case of plate production by twin roll casting or thin slab casting, the alloy utilized should have excellent formability and will be processed by hot rolling as a step in the production process. Furthermore, hot formability is an important feature of high-strength alloys with respect to their use for manufacturing parts with different shapes by methods such as hot pressing, hot stamping and the like.

  Using iron additives and other readily commercially available ingredients, 35 g of commercial purity (CP) raw material for alloy 82 representing class 2 steel was weighed according to the atomic ratios provided in Table 2. The raw material was then placed in the copper hearth of the arc melting system. The raw material was arc melted into an ingot using high purity argon as a shielding gas. The ingot was inverted several times and remelted to ensure uniformity. The resulting ingot is then placed in a PVC chamber, melted using RF induction, and then on a copper mold designed to cast a 3 × 4 inch plate having a thickness of 1.8 mm. Was released.

  The resulting plate from alloy 82 was subjected to a HIP cycle at 1150 ° C. using an American Isostatic Press Model 645 instrument with a molybdenum furnace having a 5 inch height and a 4 inch diameter furnace chamber size. The plate was heated to the target temperature at 10 ° C./min and exposed to an isostatic pressure of 30 ksi for 1 hour. A heat treatment at 850 ° C. for 1 hour was applied after the HIP cycle. Tensile specimens having a gauge length of 12 mm and a width of 3 mm were cut from the treated plate.

  Tensile measurements were made with the test parameters listed in Table 3 at the temperatures specified in Table 4. The nanosteel R & D specimen shape (shown in FIG. 11) was modified by enlarging the grip to fit the required pinhole for high temperature tensile testing. The modified grip of the sample is 9.5 mm (3/8 ″). In Table 5, a summary of the tensile test results including total tensile elongation (strain), yield stress and ultimate tensile strength is from Alloy 82. The room temperature tensile property range for the same alloy after the same treatment is given for comparison: As can be seen, the ductility in the high strength alloy is twice as high at 700 ° C., 800 It reaches up to 92% when tested at 0 ° C., demonstrating the high warm formability of the alloy.The hot temperature ductility of the alloy is strongly dependent on the alloy chemistry, thermomechanical processing parameters and test temperature.

Example # 2: Warm formability of class 2 non-stainless alloy Using iron additives and other readily commercially available ingredients, 35 g of commercial purity (CP) raw material for alloy 213 representing class 2 steel is Were weighed according to the atomic ratios provided in Table 2. The raw material was then placed in the copper hearth of the arc melting system. The raw material was arc melted into an ingot using high purity argon as a shielding gas. The ingot was inverted several times and remelted to ensure uniformity. The resulting ingot is then placed in a PVC chamber, melted using RF induction, and then on a copper mold designed to cast a 3 × 4 inch plate having a thickness of 1.8 mm. Was released.

  The resulting plate from alloy 213 was subjected to a HIP cycle at 1125 ° C. using an American Isostatic Press Model 645 instrument with a molybdenum furnace having a 5 inch height and a 4 inch diameter furnace chamber size. The plate was heated to the target temperature at 10 ° C./min and exposed to an isostatic pressure of 30 ksi for 1 hour. Tensile specimens with nanosteel R & D specimen geometry (FIG. 11) were cut from the treated plates.

  Tensile measurements were made with the test parameters listed in Table 6 at the temperatures specified in Table 7. In Table 8, a summary of the tensile test results including total tensile elongation (strain), yield stress and ultimate tensile strength is shown for the treated plate from alloy 213. The room temperature tensile property range for the same alloy after the same treatment is given for comparison. As can be seen, this alloy exhibits a high ductility of up to 74% when processed at 700 ° C., demonstrating high warm formability. The temperature dependence of tensile elongation and yield stress is shown on FIG. The hot temperature ductility of the alloy is strongly dependent on the alloy chemistry, thermomechanical processing parameters and test temperature.

Case # 3: Warm formability of class 3 alloys A survey was conducted to evaluate the warm formability of the alloys described herein at elevated temperatures. In the case of plate production by twin roll casting or thin slab casting, the alloy utilized should have excellent formability and will be processed by hot rolling as a step in the production process. Furthermore, hot formability is an important feature of high-strength alloys with respect to their use for manufacturing parts with different shapes by methods such as hot pressing, hot stamping and the like.

  Using high purity elements, 35 g of alloy raw material of alloy 36 representing class 3 steel was weighed according to the atomic ratios provided in Table 2. The raw material was then placed in the copper hearth of the arc melting system. The raw material was arc melted into an ingot using high purity argon as a shielding gas. The ingot was inverted several times and remelted to ensure uniformity. The resulting ingot is then placed in a PVC chamber, melted using RF induction, and then on a copper mold designed to cast a 3 × 4 inch plate having a thickness of 1.8 mm. Was released.

  The resulting plate from Alloy 36 was subjected to a HIP cycle at 1100 ° C. using an American Isostatic Press Model 645 instrument with a molybdenum furnace having a 5 inch height and a 4 inch diameter furnace chamber size. The plate was heated to the target temperature at 10 ° C./min and exposed to an isostatic pressure of 30 ksi for 1 hour. A heat treatment at 850 ° C. for 1 hour was applied after the HIP cycle. Tensile specimens with nanosteel R & D specimen geometry (FIG. 11) were cut from the treated plates.

Tensile measurements were made at 700 ° C. with a strain rate of 0.001 s ⁻¹ . In Table 9, a summary of the tensile test results including total tensile elongation (strain), yield stress and ultimate tensile strength is shown for the treated plate from alloy 36. The room temperature tensile property range for the same alloy after the same treatment is given for comparison. As can be seen, high strength alloys with ultimate strength up to 1650 MPa at room temperature show high ductility up to 88.5% when tested at 700 ° C., demonstrating high warm formability. The hot temperature ductility of the alloy is strongly dependent on the alloy chemistry, thermomechanical processing parameters and test temperature. An example of a specimen tested is shown in FIG.

Example # 4: Warm formability of commercial sheet from class 2 alloy Alloy 82 is by thin strip casting by in-line hot rolling performed at 1050 ° C. to ˜9% reduction. Used for the manufacture of commercial sheets. The conditions of the sheet material are not optimized (in-line rolling reduction and partial transformation to a nanomodal structure due to low temperature). Tensile specimens (FIG. 11) with nanosteel R & D specimen shapes were cut from the manufactured sheets. Tensile measurements were made with the test parameters listed in Table 10 at the temperatures specified in Table 11. In Table 12, a summary of tensile test results including total tensile elongation (strain), yield stress and ultimate tensile strength is shown for manufactured sheets from alloy 82. The temperature dependence of tensile elongation and strength properties is shown in FIG. As can be seen, ductility up to 30% can be achieved at 700 ° C., despite only partial transformation to nanomodal structure in in-line hot rolling. Higher warm formability is expected in sheets with complete transformation.

Claims (16)

Fe of 48.0 atomic% to 81.0 atomic%, B of 2.0 atomic% to 8.0 atomic%, Si of 4.0 atomic% to 14.0 atomic%, and Cu, Mn or Ni Supplying a metal alloy containing at least one of them, wherein Cu is present at 0.1 atomic percent to 6.0 atomic percent, and Mn is present at 0.1 atomic percent to 21.0 atomic percent And Ni is present at 0.1 atomic percent to 16.0 atomic percent;
Melting and solidifying the alloy to form a matrix particle size of 500 nm to 20,000 nm and a boride particle size of 25 nm to 500 nm;
Subjecting the alloy to mechanical stress and / or heating to form at least one of the following:
(A) 500 nm to 20,000 nm matrix particle size, 25 nm to 500 nm boride particles, 1 nm to 200 nm precipitation particle size, provided that the alloy has a yield strength of 400 MPa to 1300 MPa, a tensile strength of 700 MPa to 1400 MPa, and 10 % To 50% tensile elongation;
(B) A fine matrix particle size of 100 nm to 2000 nm, a precipitation particle size of 1 nm to 200 nm, a boride particle size of 200 nm to 2500 nm, provided that the alloy has a yield strength of 300 MPa to 800 MPa.   The method of claim 1, wherein the alloy of (a) is heated at a temperature of 200 ° C to 850 ° C for a period of time up to 1 hour and no eutectoid transformation is present upon cooling.   The method of claim 1, wherein the alloy of (b) is heated at a temperature of 200 ° C to 850 ° C for a period of time up to 1 hour, and no eutectoid transformation is present upon cooling.   The method of claim 2, wherein the alloy is formed into a selected shape.   The method of claim 3, wherein the alloy is formed into a selected shape.   The alloy having the fine matrix particle size (b) is exposed to stress exceeding the yield strength of 300 MPa to 800 MPa, the fine matrix particle size remains from 100 nm to 2000 nm, and the boride particle size is from 200 nm. The precipitation particle size remains at 1 nm to 200 nm, and the alloy exhibits a yield strength of 400 MPa to 1700 MPa, a tensile strength of 800 MPa to 1800 MPa, and an elongation of 5% to 40%. The method described in 1.   The method of claim 6, wherein the alloy is heated at a temperature of 200 ° C. to 850 ° C. for a period of 1 hour, and no eutectoid transformation is present upon cooling.   The method of claim 7, wherein the alloy is formed into a selected shape.   The method of claim 1, comprising Cr at a level up to 32 atomic%.   The method of claim 1 comprising C, Al, Ti, V, Nb, Mo, Zr, W or Pd at a level up to 10 atomic%. (A) Fe from 48.0 atomic% to 81.0 atomic%, B from 2.0 atomic% to 8.0 atomic%, Si from 4.0 atomic% to 14.0 atomic%, and Cu, Mn Or supplying a metal alloy containing at least one of Ni, wherein Cu is present at 0.1 atomic percent to 6.0 atomic percent, and Mn is 0.1 atomic percent to 21.0 atomic percent And Ni is present at 0.1 atomic percent to 16.0 atomic percent;
(B) melting and solidifying the alloy to provide a dendritic morphology and a matrix particle size of 500 nm to 20,000 nm and a boride particle size of 100 nm to 2500 nm;
(C) heat treating the alloy to form a lath structure containing boride grains of 100 nm to 2500 nm and grains of 100 nm to 10,000 nm, wherein the alloy has a yield strength of 300 MPa to 1400 MPa, 350 MPa to 1600 MPa Having a tensile strength of and an elongation of 0% to 12%;
(D) Heat treating the alloy after step (c) to combine 100 nm to 25000 nm boride grains and 1.0 nm to 100 nm precipitate grains, 100 nm to 10,000 nm thickness, 0.1 micrometer Forming a lamellar grain having a length of from 0.5 to 5.0 micrometers and a width of from 100 nm to 1000 nm, wherein the alloy exhibits a yield strength of 350 MPa to 1400 MPa, comprising:
(E) A method wherein the alloy is heated at a temperature of 200 ° C. to 850 ° C. for a period of up to 1 hour, and no eutectoid transformation is present upon cooling.   The alloy formed in step (d) is pressured prior to step (e) to form an alloy having 100 nm to 5000 nm grains, 100 nm to 2500 nm boride grains, 1 nm to 100 nm precipitate grains; The method of claim 11, wherein the alloy has a yield strength of 500 MPa to 1800 MPa, a tensile strength of 1000 MPa to 2000 MPa, and an elongation of 0.5% to 15%.   The method of claim 11, wherein the alloy is formed into a selected shape.   The method of claim 12, wherein the alloy is formed into a selected shape.   12. The method of claim 11, comprising Cr at a level up to 32 atomic%.   12. The method of claim 11 comprising C, Al, Ti, V, Nb, Mo, Zr, W or Pd at a level up to 10 atomic%.