US20150114587A1 - Metal Steel Production by Slab Casting - Google Patents

Metal Steel Production by Slab Casting Download PDF

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Publication number
US20150114587A1
US20150114587A1 US14/525,859 US201414525859A US2015114587A1 US 20150114587 A1 US20150114587 A1 US 20150114587A1 US 201414525859 A US201414525859 A US 201414525859A US 2015114587 A1 US2015114587 A1 US 2015114587A1
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alloy
thickness
mpa
hot rolling
cast
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Inventor
Daniel James Branagan
Grant G. JUSTICE
Andrew T. BALL
Jason K. Walleser
Brian E. MEACHAM
Kurtis CLARK
Longzhou MA
Igor YAKUBTSOV
Scott LARISH
Sheng Cheng
Taylor L. GIDDENS
Andrew E. Frerichs
Alla V. Sergueeva
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Nanosteel Co Inc
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Nanosteel Co Inc
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Priority to US14/525,859 priority Critical patent/US20150114587A1/en
Priority to US14/616,296 priority patent/US9074273B2/en
Assigned to THE NANOSTEEL COMPANY, INC. reassignment THE NANOSTEEL COMPANY, INC. ASSIGNMENT OF ASSIGNORS INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: MA, LONGZHOU, YAKUBTSOV, IGOR, FRERICHS, ANDREW E., BALL, ANDREW T., BRANAGAN, DANIEL JAMES, CHENG, SHENG, CLARK, KURTIS, GIDDENS, TAYLOR L., JUSTICE, GRANT G., LARISH, SCOTT, MEACHAM, BRIAN E., SERGUEEVA, ALLA V., WALLESER, JASON K.
Publication of US20150114587A1 publication Critical patent/US20150114587A1/en
Assigned to HORIZON TECHNOLOGY FINANCE CORPORATION reassignment HORIZON TECHNOLOGY FINANCE CORPORATION SECURITY INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: THE NANOSTEEL COMPANY, INC.
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/001—Continuous casting of metals, i.e. casting in indefinite lengths of specific alloys
    • B22D11/002—Stainless steels
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/006—Continuous casting of metals, i.e. casting in indefinite lengths of tubes
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/021—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular fabrication or treatment of ingot or slab
    • C21D8/0215—Rapid solidification; Thin strip casting
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
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    • C22C38/004—Very low carbon steels, i.e. having a carbon content of less than 0,01%
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    • C—CHEMISTRY; METALLURGY
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    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
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    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/56—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.7% by weight of carbon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/001—Continuous casting of metals, i.e. casting in indefinite lengths of specific alloys
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/04—Continuous casting of metals, i.e. casting in indefinite lengths into open-ended moulds
    • B22D11/041—Continuous casting of metals, i.e. casting in indefinite lengths into open-ended moulds for vertical casting
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/06—Continuous casting of metals, i.e. casting in indefinite lengths into moulds with travelling walls, e.g. with rolls, plates, belts, caterpillars
    • B22D11/0622—Continuous casting of metals, i.e. casting in indefinite lengths into moulds with travelling walls, e.g. with rolls, plates, belts, caterpillars formed by two casting wheels
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/12—Accessories for subsequent treating or working cast stock in situ
    • B22D11/1206—Accessories for subsequent treating or working cast stock in situ for plastic shaping of strands
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/12—Accessories for subsequent treating or working cast stock in situ
    • B22D11/128—Accessories for subsequent treating or working cast stock in situ for removing
    • B22D11/1282—Vertical casting and curving the cast stock to the horizontal
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/001—Heat treatment of ferrous alloys containing Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni

Definitions

  • This application deals with metal alloys and methods of processing with application to slab casting methods with post processing steps towards sheet production. These metals provide unique structures and exhibit advanced property combinations of high strength and/or high ductility.
  • LSS Low Strength Steels
  • HSS High-Strength Steels
  • Advanced High-Strength Steels (AHSS) steels may be understood herein as having 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% to 55%, 10% to 45% and 4% to 30%, respectively.
  • Continuous casting also called strand casting, is the process whereby molten metal is solidified into a “semifinished” billet, bloom, or slab for subsequent rolling in the finishing mills.
  • steel Prior to the introduction of continuous casting in the 1950s, steel was poured into stationary molds to form ingots. Since then, “continuous casting” has evolved to achieve improved yield, quality, productivity and cost efficiency. It allows lower-cost production of metal sections with better quality, due to the inherently lower costs of continuous, standardized production of a product, as well as providing increased control over the process through automation. This process is used most frequently to cast steel (in terms of tonnage cast). Continuous casting of slabs with either in-line hot rolling mill or subsequent separate hot rolling is important post processing steps to produce coils of sheet.
  • Thick slabs are typically cast from 150 to 500 mm thick and then allowed to cool to room temperature. Subsequent hot rolling of the slabs after preheating in tunnel furnaces is done is several stages through both roughing and hot rolling mills to get down to thicknesses typically from 2 to 10 mm in thickness. Thin slab castings starts with an as-cast thickness of 20 to 150 mm and then is usually followed through in-line hot rolling in a number of steps in sequence to get down to thicknesses typically from 2 to 10 mm. There are many variations of this technique such as casting at thicknesses of 100 to 300 mm to produce intermediate thickness slabs which are subsequently hot rolled.
  • casting processes including single and double belt casting processes which produce as-cast thickness in the range of 5 to 100 mm in thickness and which are usually in-line hot rolled to reduce the gauge thickness to targeted levels for coil production.
  • forming of parts from sheet materials from coils is accomplished through many processes including bending, hot and cold press forming, drawing, or further shape rolling.
  • the present disclosure is directed at alloys and their associated methods of production.
  • the method comprises:
  • the alloy in step (c) may undergo one of the following additional steps: (1) stressing above the alloy's yield strength of 200 MPa to 1000 MPa and providing a resulting alloy that indicates a yield strength of 200 MPa to 1650 MPa, tensile strength of 400 MPa to 1825 MPa, and an elongation of 2.4% to 78.1%; or (2) heat treating the alloy to a temperature of 700° C. to 1200° C. to form an alloy having one of the following: matrix grains of 50 nm to 50000 nm; boride grains of 20 nm to 10000 nm (optional—not required); or precipitation grains with size of 1 nm to 200 nm.
  • Such alloy with such morphology after heat treatment may then be stressed above its yield strength to form an alloy having yield strength of 200 MPa to 1650 MPa, tensile strength of 400 MPa to 1825 MPa and an elongation of 2.4% to 78.1%.
  • the alloys of present disclosure have application to continuous casting processes including belt casting, thin strip/twin roll casting, thin slab casting and thick slab casting.
  • the alloys find particular application in vehicles, such as vehicle frames, drill collars, drill pipe, pipe casing, tool joint, wellhead, compressed gas storage tanks or liquefied natural gas canisters.
  • FIG. 1 illustrates a continuous slab casting process flow diagram
  • FIG. 2 illustrates an example thin slab casting process flow diagram showing steel sheet production steps.
  • FIG. 3 illustrates a hot (cold) rolling process
  • FIG. 4 illustrates the formation of Class 1 steel alloys.
  • FIG. 5 illustrates a model stress-strain curve corresponding to Class 1 alloy behavior.
  • FIG. 6 illustrates the formation of Class 2 steel alloys.
  • FIG. 7 illustrates a model stress-strain curve corresponding to Class 2 alloy behavior.
  • FIG. 8 illustrates structures and mechanisms in the alloys herein applicable to sheet production with the identification of the Mechanism #0 (Dynamic Nanophase Refinement) which is preferably applicable to the Modal Structure (Structure #1) that is formed at thicknesses greater than or equal to 2.0 mm or at cooling rates of less than or equal to 250 K/s.
  • Mechanism #0 Dynamic Nanophase Refinement
  • Structure #1 Modal Structure
  • FIG. 9 illustrates the as-cast plate of Alloy 2 with thickness of 50 mm.
  • FIG. 10 illustrates tensile properties of the plates from Alloy 1, Alloy 8 and Alloy 16 in as-cast and heat treated states.
  • FIG. 11 illustrates SEM backscattered electron images of microstructure in the Alloy 1 plates cast at 50 mm thickness (a) before and (b) after heat treatment at 1150° C. for 120 min.
  • FIG. 12 illustrates SEM backscattered electron images of microstructure in the Alloy 8 plates cast at 50 mm thickness (a) before and (b) after heat treatment at 1100° C. for 120 min.
  • FIG. 13 illustrates SEM backscattered electron images of microstructure in the Alloy 16 plates cast at 50 mm thickness (a) before and (b) after heat treatment at 1150° C. for120 min.
  • FIG. 14 illustrates tensile properties of (a) Alloy 58 and (b) Alloy 59 in as-HIPed state as a function of cast plate thickness.
  • FIG. 15 illustrates SEM backscattered electron images of microstructure in the Alloy 59 plate cast at 1.8 mm thickness: (a) as-cast and (b) after HIP.
  • FIG. 16 illustrates SEM backscattered electron images of microstructure in the Alloy 59 plate cast at 10 mm thickness (a) as-cast and (b) after HIP.
  • FIG. 17 illustrates SEM backscattered electron images of microstructure in the Alloy 59 plate cast at 20 mm thickness (a) as-cast and (b) after HIP.
  • FIG. 18 illustrates tensile properties of (a) Alloy 58 and (b) Alloy 59 after HIP cycle and heat treatment as a function of cast thickness.
  • FIG. 19 illustrates a 20 mm thick plate from Alloy 1 before hot rolling (Bottom) and after hot rolling (Top).
  • FIG. 20 illustrates tensile properties of (a) Alloy 1 and (b) Alloy 2 before and after hot rolling as a function of cast thickness.
  • FIG. 21 illustrates backscattered SEM images of microstructure in Alloy 1 plate with as-cast thickness of 5 mm after hot rolling with 75.7% reduction in (a) outer layer region and (b) central layer region.
  • FIG. 22 illustrates backscattered SEM images of microstructure in Alloy 1 plate with as-cast thickness of 10 mm after hot rolling with 88.5% reduction in (a) outer layer region and (b) central layer region.
  • FIG. 23 illustrates backscattered SEM images of microstructure in Alloy 1 plate with as-cast thickness of 20 mm after hot rolling with 83.3% reduction in (a) outer layer region and (b) central layer region.
  • FIG. 24 illustrates tensile properties of the sheet from (a) Alloy 1 and (b) Alloy 2 after hot rolling, cold rolling and heat treatment with different parameters.
  • FIG. 25 illustrates backscattered SEM images of microstructure in Alloy 1 plate with as-cast thickness of 50 mm after hot rolling with 96% reduction in (a) outer layer region and (b) central layer region.
  • FIG. 26 illustrates backscattered SEM images of microstructure in Alloy 2 plate with as-cast thickness of 50 mm after hot rolling with 96% reduction in (a) outer layer region and (b) central layer region.
  • FIG. 27 illustrates tensile properties of post-processed sheet from (a) Alloy 1 and (b) Alloy 2 at different steps of post-processing.
  • FIG. 28 illustrates tensile properties of post-processed sheet from (a) Alloy 1 and (b) Alloy 2 initially cast at different thicknesses.
  • FIG. 29 illustrates backscattered SEM images of Alloy 2 with as-cast thickness of 20 mm after hot rolling with 88% reduction: (a) outer layer region; (b) central layer region.
  • FIG. 30 illustrates backscattered SEM images of Alloy 2 20 mm thick plate sample hot rolled and heat treated at 950° C. for 6 hr: (a) outer layer region; (b) central layer region.
  • FIG. 31 illustrates tensile properties of Alloy 8 sheet produced from 50 mm thick plate by hot rolling that was heat treated at different conditions with representative stress-strain curves.
  • FIG. 32 illustrates tensile properties of Alloy 16 sheet produced from 50 mm thick plate by hot rolling that was heat treated at different conditions.
  • FIG. 33 illustrates tensile properties of Alloy 24 sheet produced from 50 mm thick plate by hot rolling that was heat treated at different conditions with representative stress-strain curves.
  • FIG. 34 illustrates bright-field TEM micrographs of microstructure in the Alloy 1 plate after hot rolling and heat treatment initially cast 50 mm thickness.
  • FIG. 35 illustrates bright-field TEM micrographs of microstructure in the hot rolling and heat treated Alloy 1 plate after tensile deformation.
  • FIG. 36 illustrates bright-field TEM micrographs of microstructure in the 50 mm thick Alloy 8 plate after hot rolling and heat treatment: (a) before and (b) after tensile deformation.
  • FIG. 37 illustrates bright-field TEM micrographs at higher magnification of microstructure in the 50 mm thick Alloy 8 plate after hot rolling and heat treatment: (a) before and (b) after tensile deformation.
  • FIG. 38 illustrates high resolution TEM micrographs of microstructure in the 50 mm thick Alloy 8 plate after hot rolling and heat treatment: (a) before and (b) after tensile deformation.
  • FIG. 39 illustrates bright-field and dark-field TEM micrographs of microstructure in the 50 mm thick Alloy 16 plate after hot rolling and heat treatment.
  • FIG. 40 illustrates bright-field and dark-field TEM micrographs of microstructure in the hot rolled and heat treated Alloy 16 plate after tensile deformation.
  • FIG. 41 illustrates tensile properties of post-processed sheet from Alloy 32 and Alloy 42 initially cast into 50 mm thick plates.
  • FIG. 42 illustrates bright-field TEM micrographs of microstructure in the 50 mm thick as-cast plate from Alloy 24.
  • FIG. 43 illustrates bright-field TEM micrographs of microstructure in the Alloy 24 plate after hot rolling from 50 to 2 mm thickness.
  • FIG. 44 illustrates schematic of the cross section through the center of the cast plate showing the shrinkage funnel and the locations from which samples for chemical analysis were taken.
  • FIG. 45 illustrates alloying element content in tested locations at the top (Area A) and bottom (Area B) of the cast plate for the four alloys identified.
  • FIG. 46 illustrates comparison of stress-strain curves of new steel sheet types with existing Dual Phase (DP) steels.
  • FIG. 47 illustrates comparison of stress-strain curves of new steel sheet types with existing Complex Phase (CP) steels.
  • FIG. 48 illustrates comparison of stress-strain curves of new steel sheet types with existing Transformation Induced Plasticity (TRIP) steels.
  • TRIP Transformation Induced Plasticity
  • FIG. 49 illustrates comparison of stress-strain curves of new steel sheet types with existing Martensitic (MS) steels.
  • FIG. 51 illustrates tensile properties of selected alloys cast at 50 mm thickness as compared to that for the same alloys cast at 3.3 mm thickness.
  • FIG. 52 illustrates an example stress strain curve of boron-free Alloy 63 in hot rolled state.
  • FIG. 53 Backscattered electron images of microstructure in the Alloy 65 cast at 50 mm thickness: (a) as-cast; (b) after hot rolling at 1250° C.; (c) after cold rolling to 1.2 mm thickness.
  • a slab is a length of metal that is rectangular in cross-section.
  • Slabs can be produced directly by continuous casting and are usually further processed via different processes (hot/cold rolling, skin rolling, batch heat treatment, continuous heat treatment, etc.). Common final products include sheet metal, plates, strip metal, pipes, and tubes.
  • Thick slab casting is the process whereby molten metal is solidified into a “semifinished” slab for subsequent rolling in the finishing mills.
  • molten steel flows from a ladle, through a tundish into the mold. Once in the mold, the molten steel freezes against the water-cooled copper mold walls to form a solid shell.
  • Drive rolls lower in the machine continuously withdraw the shell from the mold at a rate or “casting speed” that matches the flow of incoming metal, so the process ideally runs in steady state.
  • the solidifying steel shell acts as a container to support the remaining liquid. Rolls support the steel to minimize bulging due to the ferrostatic pressure.
  • Water and air mist sprays cool the surface of the strand between rolls to maintain its surface temperature until the molten core is solid.
  • the strand can be torch cut into slabs with typical thickness of 150 to 500 mm.
  • hot rolling may be done in both roughing mills which are often reversible allowing multiple passes and with finishing fills with typically 5 to 7 stands in series. After hot rolling, the resulting sheet thickness is typically in the range of 2 to 5 mm. Further gauge reduction would occur normally through subsequent cold rolling.
  • FIG. 2 A schematic of the thin slab casting process is shown in FIG. 2 .
  • the thin slab casting process can be separated into three stages.
  • Stage 1 the liquid steel is both cast and rolled in an almost simultaneous fashion.
  • the solidification process begins by forcing the liquid melt through a copper or copper alloy mold to produce initial thickness typically from 50 to 110 mm in thickness but this can be varied (i.e. 20 to 150 mm) based on liquid metal processability and production speed. Almost immediately after leaving the mold and while the inner core of the steel sheet is still liquid, the sheet undergoes reduction using a multistep rolling stand which reduces the thickness significantly down to 10 mm depending on final sheet thickness targets.
  • Stage 2 the steel sheet is heated by going through one or two induction furnaces and during this stage the temperature profile and the metallurgical structure is homogenized.
  • the sheet is further rolled to the final gage thickness target which may be in the 0.5 to 15 mm thickness range.
  • the gauge reduction will be done in 5 to 7 steps as the sheet is reduced through 5 to 7 mills in series.
  • the strip is cooled on a run-out table to control the development of the final microstructure of the sheet prior to coiling into a steel roll.
  • Hot rolled steel is formed to shape while it is red-hot then allowed to cool.
  • Flat rolling is the most basic form of rolling with the starting and ending material having a rectangular cross-section.
  • the schematic illustration of a rolling process for metal sheets is presented in FIG. 3 .
  • Hot rolling is a part of sheet production in order to reduce sheet thickness towards targeted values by utilizing the enhanced ductility of sheet metal at elevated temperature when high level of rolling reduction can be achieved.
  • Hot rolling can be a part of casting process when one (Thin Strip casting) or multiple (Thin Slab Casting) stands are built-in in-line. In a case of Thick (Traditional) Slab Casting, the slab is first reheated in a tunnel furnace and then moves through a series of mill stands ( FIG. 3 ).
  • hot rolling is a part of post-processing on separate Hot Rolling Mill Production Lines is also applied. Since red-hot steel contracts as it cools, the surface of the metal is slightly rough and the thickness may vary a few thousandths of an inch. Commonly, cold rolling is a following step to improve quality in the final sheet product.
  • Cold rolled steel is made by passing cold steel material through heavy rollers which compress the metal to its final shape and dimension. It is a common step of post-processing during sheet production when different cold rolling mills can be utilized depending on material properties, cold rolling objective and targeted parameters.
  • sheet material undergoes cold rolling its strength, hardness as well as the elastic limit increase.
  • the ductility of the metal sheet decreases due to strain hardening thus making the metal more brittle.
  • the metal must be annealed/heated from time to time between passes during the rolling operation to remove the undesirable effects of cold deformation and to increase the formability of the metal. Thus obtaining large thickness reduction can be time and cost consuming.
  • multi-stand cold rolling mills with in-line annealing are utilized wherein the sheet is affected by elevated temperature for a short period of time (usually 2 to 5 min) by induction heating while it moves along the rolling line.
  • Cold rolling allows a much more precise dimensional accuracy and final sheet products have a smoother surface (better surface finish) than those from hot rolling.
  • annealing of steel sheet products is usually implemented.
  • annealing of steel sheet products is performed in two ways at a commercial scale: batch annealing or continuous annealing.
  • batch annealing process massive coils of the sheet slowly heat and cool in furnaces with a controlled atmosphere.
  • the annealing time can be from several hours to several days. Due to the large mass of the coils which may be typically 5 to 25 ton in size, the inside and outside parts of the coils will experience different thermal histories in a batch annealing furnace which can lead to differences in resulting properties.
  • a continuous annealing process uncoiled steel sheets pass through heating and cooling equipment for several minutes.
  • the heating equipment is usually a two-stage furnace.
  • the first stage is high temperature heat treatment which provides recrystallization of microstructure.
  • the second stage is low temperature heat treatment and it offers artificial ageing of microstructure.
  • a proper combination of the two stages of overall heat treatment during continuous annealing provides the target mechanical properties.
  • the advantages of continuous annealing over conventional batch annealing are the following: improved product uniformity; surface cleanliness and shape; ability to produce a wide range of steel grades.
  • the steel alloys herein are such that they are initially capable of formation of what is described herein as Class 1 or Class 2 Steel which are preferably crystalline (non-glassy) with identifiable crystalline grain size and morphology.
  • Class 1 or Class 2 Steel which are preferably crystalline (non-glassy) with identifiable crystalline grain size and morphology.
  • the present disclosure focuses upon improvements to the Class 2 Steel and the discussion below regarding Class 1 is intended to provide initial context.
  • Class 1 Steel herein is illustrated in FIG. 4 .
  • a modal structure is initially formed which modal structure is the result of starting with a liquid melt of the alloy and solidifying by cooling, which provides nucleation and growth of particular phases having particular grain sizes.
  • Reference herein to modal may therefore be understood as a structure having at least two grain size distributions.
  • Grain size herein may be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy.
  • Structure #1 of the Class 1 Steel may be preferably achieved by processing through either laboratory scale procedures as shown and/or through industrial scale methods involving chill surface processing methodology such as twin roll processing, thin slab casting or thick slab casting.
  • the modal structure of Class 1 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 500 nm to 20,000 nm containing austenite and/or ferrite; (2) boride grain size of 25 nm to 5000 nm (i.e. non-metallic grains such as M 2 B 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.
  • metal boride grains have been identified as exhibiting the M 2 B stoichiometry but other stoichiometry is possible and may provide pinning including M 3 B, MB (M 1 B 1 ), M 23 B 6 , and M 7 B 3 .
  • the Modal Structure of Class 1 Steel may be deformed by thermo-mechanical processes and undergo various heat treatments, resulting in some variation in properties, but the Modal Structure may be maintained.
  • the observed stress versus strain diagram is illustrated in FIG. 5 . It is therefore observed that the modal structure undergoes what is identified as the Dynamic Nanophase Precipitation leading to a second type structure for the Class 1 Steel. Such Dynamic Nanophase Precipitation is therefore triggered when the alloy experiences a yield under stress, and it has been found that the yield strength of Class 1 Steels which undergo Dynamic Nanophase Precipitation may preferably occur at 300 MPa to 840 MPa. Accordingly, it may be appreciated that the Dynamic Nanophase Precipitation occurs due to the application of mechanical stress that exceeds such indicated yield strength.
  • the Dynamic Nanophase Precipitation itself may be understood as the formation of a further identifiable phase in the Class 1 Steel which is termed a precipitation phase with an associated grain size. That is, the result of such Dynamic Nanophase Precipitation is to form an alloy which still indicates identifiable matrix grain size of 500 nm to 20,000 nm, boride pinning grain size of 20 nm to 10000 nm, along with the formation of precipitation grains of hexagonal phases with 1.0 nm to 200 nm in size As noted above, the grain sizes therefore do not coarsen when the alloy is stressed, but does lead to the development of the precipitation grains as noted.
  • references to the hexagonal phases may be understood as a dihexagonal pyramidal class hexagonal phase with a P6 3 mc space group (#186) and/or a ditrigonal dipyramidal class with a hexagonal P6bar2C space group (#190).
  • 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 1150 MPa, with an elongation of 10 to 40%.
  • 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
  • ⁇ represents the applied stress on the material
  • ⁇ is the strain
  • 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 herein is illustrated in FIG. 6 .
  • Class 2 steel may also be formed herein from the identified alloys, which involves two new structure types after starting with Structure #1, Modal Structure, followed by two new mechanisms identified herein as Static Nanophase Refinement and Dynamic Nanophase Strengthening.
  • the structure types for Class 2 Steel are described herein as Nanomodal Structure and High Strength Nanomodal Structure. Accordingly, Class 2 Steel herein may be characterized as 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).
  • Structure #1 is initially formed in which Modal Structure is the result of starting with a liquid melt of the alloy and solidifying by cooling, which provides nucleation and growth of particular phases having particular grain sizes.
  • Grain size herein may again be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy.
  • Structure #1 of the Class 2 Steel may be preferably achieved by processing through either laboratory scale procedures as shown and/or through industrial scale methods involving chill surface processing methodology such as twin roll processing or thin slab casting.
  • the Modal Structure of Class 2 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 200 nm to 200,000 nm containing austenite and/or ferrite; (2) boride grain sizes, if present, of 10 nm to 5000 nm (i.e. non-metallic grains such as M 2 B where M is the metal and is covalently bonded to B).
  • the boride grains may also preferably be “pinning” type phases which are referenced to the feature that the matrix grains will effectively be stabilized by the pinning phases which resist coarsening at elevated temperature.
  • metal boride grains have been identified as exhibiting the M 2 B stoichiometry but other stoichiometry is possible and may provide pinning including M 3 B, MB (M 1 B 1 ), M 23 B 6 , and M 7 B 3 and which are unaffected by Mechanisms #1 or #2 noted above.
  • Reference to grain size is again to be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy.
  • Structure #1 of Class 2 steel herein includes austenite and/or ferrite along with such boride phases.
  • a stress strain curve is shown that represents the steel alloys herein which undergo a deformation behavior of Class 2 steel.
  • the Modal Structure is preferably first created (Structure #1) and then after the creation, the Modal Structure may now be uniquely refined through Mechanism #1, which is a Static Nanophase Refinement mechanism, leading to Structure #2.
  • Static Nanophase Refinement is reference to the feature that the matrix grain sizes of Structure #1 which initially fall in the range of 200 nm to 200,000 nm are reduced in size to provide Structure 2 which has matrix grain sizes that typically fall in the range of 50 nm to 5000 nm.
  • the boride pinning phase if present, can change size significantly in some alloys, while it is designed to resist matrix grain coarsening during the heat treatments. Due to the presence of these boride pinning sites, the motion of a grain boundaries leading to coarsening would be expected to be retarded by a process called Zener pinning or Zener drag. Thus, while grain growth of the matrix may be energetically favorable due to the reduction of total interfacial area, the presence of the boride pinning phase will counteract this driving force of coarsening due to the high interfacial energies of these phases.
  • Characteristic of the Static Nanophase Refinement (Mechanism #1) in Class 2 steel, if borides are present, is such that the micron scale austenite phase (gamma-Fe) which was noted as falling in the range of 200 nm to 200,000 nm is partially or completely transformed into new phases (e.g. ferrite or alpha-Fe) at elevated temperature.
  • the volume fraction of ferrite (alpha-iron) initially present in the modal structure (Structure 1) of Class 2 steel is 0 to 45%.
  • the volume fraction of ferrite (alpha-iron) in Structure #2 as a result of Static Nanophase Refinement (Mechanism #2) is typically from 20 to 80% at elevated temperature and then reverts back to austenite (gamma-iron) upon cooling to produce typically from 20 to 80% austenite at room temperature.
  • the static transformation preferably occurs during elevated temperature heat treatment and thus involves a unique refinement mechanism since grain coarsening rather than grain refinement is the conventional material response at elevated temperature.
  • 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 400 to 1825 MPa with 2.4 to 78.1% total elongation.
  • nanoscale precipitates can form during Static Nanophase Refinement and the subsequent thermal process in some of the non-stainless high-strength steels.
  • the nano-precipitates are in the range of 1 nm to 200 nm, with the majority (>50%) of these phases 10 ⁇ 20 nm in size, which are much smaller than matrix grains or the boride pinning phase formed in Structure #1 for retarding matrix grain coarsening when present.
  • the boride grains if present, are found to be in a range from 20 to 10000 nm in size.
  • 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 200 MPa to 1650 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. 200 to 1650 MPa) as applied to the Structure #2 transformation into Structure #3, allowing tunable variations to enable both the designer and end users in a variety of applications, and utilize Structure #3 in various applications such as crash management in automobile body structures.
  • a wide range e.g. 200 to 1650 MPa
  • Structure #3 may be understood as a microstructure having matrix grains sized generally from 25 nm to 2500 nm which are pinned by boride phases, which are in the range of 20 nm to 10000 nm and with precipitate phases which are in the range of 1 nm to 200 nm. Note that in the absence of boride pinning phases, the refinement may be somewhat less and/or some matrix coarsening may occur resulting in matrix grains which are sized from 25 nm to 25000 nm.
  • the initial formation of the above referenced precipitation phase with grain sizes of 1 nm to 200 nm starts at Static Nanophase Refinement and continues during Dynamic Nanophase Strengthening leading to Structure #3 formation.
  • the volume fraction of the precipitation grains with 1 nm to 200 nm in size increases in Structure #3 as compared to Structure #2 and assists with the identified strengthening mechanism. It should also be noted that in Structure #3, the level of gamma-iron is optional and may be eliminated depending on the specific alloy chemistry and austenite stability. Table 2 below provides a comparison of the structure and performance of Class 2 Steel herein:
  • metal borides e.g. metal borides (e.g. metal boride) (if present) boride) boride) Precipitation — 1 nm to 200 nm 1 nm to 200 nm Grain Size Tensile Response Actual with properties Intermediate structure; Actual with properties achieved based on transforms into Structure achieved based on formation structure type #1 #3 when undergoing of structure type #3 and yield fraction of transformation.
  • FIG. 8 A new pathway is disclosed herein as shown in FIG. 8 .
  • This figure relates to the alloys in which boride pinning phase may or may not be present. It starts with Structure #1, Modal Structure but includes additional Mechanism #0—Dynamic Nanophase Refinement leading to formation of Structure #1a—Homogenized Modal Structure ( FIG. 8 ). More specifically, Dynamic Nanophase Refinement is the application of elevated temperature (700° C.
  • the Dynamic Nanophase Refinement leading to the Homogenized Modal Structure is observed to occur in as little as 1 cycle (heating with thickness reduction) or after multiple reduction cycles of thickness (e.g. up to 25).
  • the Homogenized Modal Structure (Structure 1a in FIG. 8 ) represents an intermediate structure between the starting Modal Structure with the associated properties and characteristics defined as Structure 1 of FIG. 8 . and the fully transformed Nanomodal Structure defined as Structure 2 in FIG. 8 .
  • the transformation can be complete in as little as 1 cycle or it may take many cycles ((e.g. up to 25) to completely transform.
  • a partially transformed, intermediate structure is Structure 1a or Homogenized Modal Structure and after full transformation of the Modal Structure into NanoModal Structure, the Nanomodal structure (i.e. Structure 2) is formed. Progressive cycles lead to the creation of Structure #2 (Nanomodal Structure).
  • Structure #1a Homogenized Modal Structure
  • Structure #2 may therefore become directly Structure #2 (Nanomodal Structure) or may be heat treated and further refined through Mechanism #1 (Static Nanophase Refinement) to similarly produce Structure #2 (Nanomodal Structure).
  • Structure #2, Nanomodal Structure may then undergo Mechanism #2 (Dynamic Nanophase Strengthening) leading to the formation of Structure #3 (High Strength Nanomodal Structure).
  • Dynamic Nanophase Refinement is a mechanism providing Homogenized Modal Structure (Structure #1a) in cast alloys preferably through the entire volume/thickness that makes the alloys effectively cooling rate insensitive (as well as thickness insensitive) during the initial solidification from the liquid state that enables utilization of such production methods as thin slab or thick slab casting for sheet production.
  • Static Nanophase Refinement may not readily occur.
  • Dynamic Nanophase Refinement occurs after the alloys are subjected to deformation at elevated temperature and preferably occurs at a range from 700° C. to a temperature just below the melting point and over a range of strain rates from 10 ⁇ 6 to 10 4 s ⁇ 1 .
  • deformation may occur by hot rolling after thick slab or thin slab casting which may occur in single or multiple roughing hot rolling steps or single and/or single or multiple finishing hot rolling steps.
  • hot processing steps including but not limited to hot stamping, forging, hot pressing, hot extrusion, etc.
  • Modal Structure (Structure #1) in steel alloys herein can occur during alloy solidification at Thick Slab ( FIG. 1 ) or Thin Slab Casting (Stage 1, FIG. 2 ).
  • the Modal Structure 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 1 ⁇ 10 3 to 1 ⁇ 10 ⁇ 3 K/s.
  • Integrated hot rolling of Thick Slab ( FIG. 1 ) or Thin Slab Casting (Stage 2, FIG. 2 ) of the alloys will lead to formation of Homogenized Modal Structure (Structure #1a, FIG. 8 ) through the Dynamic Nanophase Refinement (Mechanism #0) in the cast slab with thickness of typically 150 to 500 mm in a case of Thick Slab Casting and 20 to 150 mm in a case of Thin Slab Casting.
  • the Type of the Homogenized Modal Structure (Table 1) will depend on alloy chemistry and hot rolling parameters.
  • Mechanism #1 which is the Static Nanophase Refinement with Nanomodal Structure formation (Structure #2) occurs when produced slabs with Homogenized Modal Structure (Structure #1a, FIG. 8 ) are subjected to elevated temperature exposure (from 700° C. up to the melting temperature of the alloy) during post-processing.
  • Possible methods for realization of Static Nanophase Refinement (Mechanism #1) include but not limited to in-line annealing, batch annealing, hot rolling followed by annealing towards targeted thickness, etc. Hot rolling is a typical method utilized to reduce slab thickness to the ranges of few millimeters in order to produce sheet steel for various applications. Typical thickness reduction can vary widely depending on the production method of the initial sheet. Starting thickness may vary from 3 to 500 mm and final thickness would vary from 1 mm to 20 mm.
  • Cold rolling is a widely used method for sheet production that is utilized to achieve targeted thickness for particular applications.
  • most sheet steel used for automotive industry has thickness in a range from 0.4 to 4 mm.
  • cold rolling is applied through multiple passes with intermediate annealing between passes. Typical reduction per pass is 5 to 70% depending on the material properties. The number of passes before the intermediate annealing also depends on materials properties and its level of strain hardening at cold deformation.
  • Cold rolling is also used as a final step for surface quality known as a skin pass.
  • the cold rolling will trigger Dynamic Nanophase Strengthening and the formation of the High Strength Nanomodal Structure.
  • Alloy 1 through Alloy 59 were cast into plates with thickness of 3.3 mm.
  • 35 g alloy feedstocks of the targeted alloys were weighed out according to the atomic ratios provided in Table 4.
  • 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. Individually, the ingots were disc-shaped, with a diameter of approximately 30 mm and a thickness of approximately 9.5 mm at the thickest point.
  • the resulting ingots were then placed in a pressure vacuum caster (PVC) chamber, melted using RF induction and then ejected onto a copper die designed for casting 3 by 4 inches sheets with thickness of 3.3 mm.
  • PVC pressure vacuum caster
  • Alloy 60 through Alloy 62 were cast into plates with thickness of 50 mm. These chemistries have been used for material processing through slab casting in an Indutherm VTC800V vacuum tilt casting machine. Alloys of designated compositions were weighed out in 3 kilogram charges using designated quantities of commercially-available ferroadditive powders of known composition and impurity content, and additional alloying elements as needed, according to the atomic ratios provided in Table 4 for each alloy. Alloy charges were placed in zirconia coated silica-based crucibles and loaded into the casting machine. Melting took place under vacuum using a 14 kHz RF induction coil.
  • the alloys herein that are susceptible to the transformations illustrated in FIG. 8 fall into the following groupings: (1) Fe/Cr/Ni/Mn/B/Si/Cu (alloys 1, 2, 15 to 18, 27 to 28, 35, 40, 50 to 57, 59, 62); (2) Fe/Ni/Mn/B/Si/Cu (alloys 3 to 6, 19, 29 to 30); (3) Fe/Mn/B/Si (alloys 7 to 10, 20, 25 to 26); (4) Fe/Cr/Mn/B/Si (alloys 11 to 14, 21 to 24, 37 to 39); Fe/Ni/Mn/B/Si/Cu/C (alloys 31, 36, 46 to 47, 61); (5) Fe/Cr/Ni/Mn/B/Si/Cu/C (alloys 32 to 34, 41 to 45, 49, 60); (6) Fe/Cr/Mn/B/Si/C (alloy 48); (7) Fe/C
  • the alloy composition herein would include the following four elements at the following indicated atomic percent: Fe (61.0 to 88.0 at. %); Si (0.5 to 9.0 at. %); Mn (0.9 to 19.0 at. %) and optionally B (0.0 at. % to 8.0 at. %).
  • the following elements are optional and may be present at the indicated atomic percent: Ni (0.1 to 9.0 at. %); Cr (0.1 to 19.0 at. %); Cu (0.1 to 4.0 at. %); C (0.1 to 4.0 at. %).
  • Impurities may be present include Al, Mo, Nb, S, O, N, P, W, Co, Sn, Zr, Ti, Pd and V, which may be present up to 10 atomic percent.
  • the alloys may herein also be more broadly described as Fe based alloys (greater than 60.0 atomic percent) and further including B, Si and Mn.
  • the alloys are capable of being solidified from the melt to form Modal Structure (Structure #1, FIG. 8 ), when at a thickness of greater than or equal to 2.0 mm, or which Modal Structure when formed at a cooling rate of less than or equal to 250 K/s, can preferably undergo Dynamic Nanophase Refinement which may then provide Homogenized Modal Structure (Structure #1a, FIG. 8 ). As indicated in FIG. 8 , one may then, from such Homogenized Modal Structure, ultimately form High Strength Nanomodal Structure (Structure #3) with the indicted morphology and mechanical properties.
  • Final melting temperature is >1425° C. in selected alloys. Liquidus temperature for these alloys is out of measurable range and not available (marked as “NA” in the Table 5). Variations in melting behavior may reflect a complex phase formation during 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 6 and was found to vary from 7.55 g/cm 3 to 7.89 g/cm 3 .
  • the accuracy of this technique is ⁇ 0.01 g/cm 3 .
  • the tensile specimens were cut from the hot rolled and heat treated sheets using wire electrical discharge machining (EDM).
  • EDM wire electrical discharge machining
  • 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 rigid and the top fixture moving; the load cell is attached to the top fixture.
  • Table 8 a summary of the tensile test results including, yield stress, ultimate tensile strength, and total elongation are shown for the hot rolled sheets after heat treatment.
  • 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 431 to 1612 MPa.
  • the tensile elongation varies from 2.4 to 64.7%. Yield stress is measured in a range from 212 MPa to 966 MPa.
  • Mechanism #2 Dynamic Nanophase Strengthening
  • All cast plates with initial thickness of 50 mm were subjected to hot rolling at the temperature of 1075 to 1100° C. depending on alloy solidus temperature. Rolling was done on a Fenn Model 061 single stage rolling mill, employing an in-line Lucifer EHS3GT-B18 tunnel furnace. Material was held at the hot rolling temperature for an initial dwell time of 40 minutes to ensure homogeneous temperature. After each pass on the rolling mill, the sample was returned to the tunnel furnace with a 4 minute temperature recovery hold to correct for temperature lost during the hot rolling pass. Hot rolling was conducted in two campaigns, with the first campaign achieving approximately 85% total reduction to a thickness of 6 mm.
  • Hot-rolled sheets from each alloy were then subjected to further cold rolling in multiple passes down to thickness of 1.2 mm. Rolling was done on a Fenn Model 061 single stage rolling mill. Examples of specific cold rolling parameters used for the alloys are shown in Table 10.
  • Tensile properties of the selected alloys after hot rolling with subsequent cold rolling and heat treatment at different parameters are listed in Table 13.
  • the ultimate tensile strength values may vary from 813 MPa to 1316 MPa with tensile elongation from 6.6 to 35.9%.
  • the yield stress is in a range from 274 to 815 MPa. This corresponds to Structure 2 in FIG. 8 .
  • the mechanical characteristic values in the steel alloys herein will depend on alloy chemistry and processing conditions.
  • FIG. 9 An example of cast plate from Alloy 2 with thickness of 50 mm is shown in FIG. 9 .
  • Sheet samples produced by multi-pass hot rolling of cast plates were the subject for further treatments (heat treatment, cold rolling, etc.) as described in the Case Examples herein mimicking sheet post-processing after Thin Slab Production depending on property and performance requirements for different applications.
  • Close modeling of the Slab Casting process and post-processing methods allow prediction of structural development in the steel alloys herein at each step of the processing and identifies the mechanisms which will lead to production of sheet steel with advanced property combinations.
  • Tensile specimens were cut from the as-cast and heat treated plates using a Brother HS-3100 wire electrical discharge machining (EDM). The tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. A video extensometer was utilized for strain measurements.
  • FIGS. 12 through 14 demonstrate SEM images of the microstructure in all three alloys before and after heat treatment.
  • Modal Structure (Structure #1) is present in as-cast plates from all three alloys with boride phase located between matrix grains and along the matrix grain boundaries.
  • Static Nanophase Refinement Mechanism #1, FIG. 8
  • the microstructure appears to remain coarse and additionally only partial spheroidization of the boundary boride phase can be seen after heat treatment with localization along prior dendrite boundaries.
  • heat treatment of the plates directly after solidification does not provide refinement and structural homogenization necessary to achieve the properties when alloys are cast at large thicknesses, resulting in relatively poor properties.
  • Static Nanophase Refinement occurring through elevated temperature heat treatment is found to be relatively ineffective in samples cast at high thickness/reduced cooling rates.
  • the range where Static Nanophase Refinement will not be effective will be dependent on the specific alloy chemistry and size of the dendrites in the Modal Structure but generally occurs at casting thickness greater than or equal to 2.0 mm and cooling rates less than or equal to 250 K/s.
  • Plate casting with different thicknesses in a range from 1.8 mm to 20 mm was done for the Alloy 58 and Alloy 59 listed in Table 4.
  • Thin plates with as-cast thickness of 1.8 mm were cast in a Pressure Vacuum Caster (PVC).
  • PVC Pressure Vacuum Caster
  • charges of 35 g were weighed out according to the atomic ratios provided in Table 4.
  • the feedstock material was then placed into the copper hearth of an arc-melting system.
  • the feedstock was arc-melted into an ingot using high purity argon as a shielding gas.
  • the ingots were flipped several times and re-melted to ensure homogeneity.
  • the ingots were disc-shaped, with a diameter of ⁇ 30 mm and a thickness of ⁇ 9.5 mm at the thickest point.
  • the resulting ingots were then placed in a PVC chamber, melted using RF induction and then ejected into a copper die designed for casting 3 by 4 inches plates with thickness of 1.8 mm.
  • Casting of plates with thickness from 5 to 20 mm was done by using an Indutherm VTC 800 V Tilt Vacuum Caster. Using commercial purity feedstock, charges of different masses were weighed out for particular alloys according to the atomic ratios provided in Table 4. The charges were then placed into the crucible of the caster. The feedstock was melted using RF induction and then poured into a copper die designed for casting plates with dimensions described in Table 16.
  • HIP cycle was used as in-situ heat treatment and a method to remove some of the casting defects to mimic hot rolling step at slab casting. HIP cycle parameters are listed in Table 17. After HIP cycle, the plates from both alloys were heat treated in a box furnace at 900° C. for 1 hr.
  • the tensile specimens were cut from the plates in as-HIPed state as well as after HIP cycle and heat treatment 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 rigid and the top fixture moving with the load cell attached to the top fixture.
  • samples in the as-cast, HIPed and heat treated states were examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.
  • SEM specimens the cross-sections of the plate samples were cut and ground by SiC paper and then polished progressively with diamond media paste down to 1 ⁇ m grit. The final polishing was done with 0.02 ⁇ m grit SiO 2 solution.
  • FIG. 15 through FIG. 17 Examples of microstructures in the plates for Alloy 59 in the as-cast state and after HIP cycle are shown in FIG. 15 through FIG. 17 .
  • Modal Structure (Structure #1) can be observed in the plates in as-cast condition ( FIG. 15 a , FIG. 16 a , FIG. 17 a ) with increasing dendrite size as a function of cast plate thickness.
  • the Modal Structure may have partially transformed into Nanomodal Structure (Structure #2) through Static Nanophase Refinement (Mechanism #1) but the structure appears coarse (note individual grain size beyond SEM resolution). But, as it can be seen in all cases ( FIG. 15 b , FIG. 16 b , FIG.
  • boride phases are preferably aligned along primary dendrites formed at solidification.
  • Significantly smaller dendrites in the case of casting at 1.8 mm thickness results in more homogeneous distribution of borides leading to better properties as compared to that in cast plates with larger thicknesses ( FIG. 15 b ).
  • Additional heat treatment after HIP cycle results in property improvement in all plated with more pronounced effect in 1.8 mm thick plates from both alloys ( FIG. 18 ). In the samples cast at greater thickness (i.e. 5 to 20 mm), the improvement in properties are minimal.
  • Plates with different thicknesses in a range from 5 mm to 20 mm were cast from Alloy 1 and Alloy 2 using an Indutherm VTC 800 V Tilt Vacuum Caster.
  • charges of different masses were weighed out for particular alloys according to the atomic ratios provided in Table 4. The charges were then placed into the crucible of the caster. The feedstock was melted using RF induction and then poured into a copper die designed for casting plates with dimensions described in Table 15.
  • Each plate from each alloy was subjected to Hot Rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace. The plates were placed in a furnace pre-heated to 1140° C.
  • Hot rolling reduction value for each plate for both Alloys is provided in Table 18.
  • Plate casting with 50 mm thickness from Alloy 1 and Alloy 2 was done using an Indutherm VTC 800 V Tilt Vacuum Caster in order to mimic the Stage 1 of the Thin Slab Process ( FIG. 2 ).
  • Indutherm VTC 800 V Tilt Vacuum Caster in order to mimic the Stage 1 of the Thin Slab Process ( FIG. 2 ).
  • charges of different masses were weighed out for Alloy 1 and Alloy 2 according to the atomic ratios provided in Table 4. The charges were then placed into the crucible of the caster.
  • the feedstock was melted using RF induction and then poured into a copper die designed for casting plates with 50 mm thickness.
  • the plates from each alloy were subjected to Hot Rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace.
  • the plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling.
  • the plates were then repeatedly rolled at between 10% and 25% reduction per pass down to 3.5 mm thickness mimicking multi-stand hot rolling at Stage 2 during the Thin Slab Process ( FIG. 2 ) or hot rolling step at Thick Slab Casting ( FIG. 1 ).
  • the plates were placed in the furnace for 1 to 2 min between rolling steps to allow them to partially return to temperature for the next rolling pass. If the plates became too long to fit in the furnace they were cooled, cut to a shorter length, then reheated in the furnace for 60 minutes before they were rolled again towards the targeted gauge thickness. Total reduction of 93% was achieved for both alloys.
  • Hot rolled sheets were heat treatment at different parameters listed in Table 19.
  • Tensile specimens were cut from the rolled and heat treated sheets from Alloy 1 and Alloy 2 using a Brother HS-3100 wire electrical discharge machining (EDM). The tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. A non-contact video extensometer was utilized for strain measurements.
  • Plate casting with 50 mm thickness from Alloy 1 and Alloy 2 was done using an Indutherm VTC 800 V Tilt Vacuum Caster in order to mimic the Stage 1 of the Thin Slab Process ( FIG. 2 ).
  • Indutherm VTC 800 V Tilt Vacuum Caster in order to mimic the Stage 1 of the Thin Slab Process ( FIG. 2 ).
  • charges of different masses were weighed out for Alloy 1 and Alloy 2 according to the atomic ratios provided in Table 4. The charges were then placed into the crucible of the caster.
  • the feedstock was melted using RF induction and then poured into a copper die designed for casting plates with 50 mm thickness.
  • the plates from each alloy were subjected to hot rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace.
  • the plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling.
  • the plates were then repeatedly rolled at between 10% and 25% reduction per pass down to 3.5 mm thickness mimicking multi-stand hot rolling at Stage 2 during the Thin Slab Process ( FIG. 2 ) or hot rolling step at Thick Slab Casting ( FIG. 1 ).
  • the plates were placed in the furnace for 1 to 2 min between rolling steps to allow them to return to temperature. If the plates became too long to fit in the furnace they were cooled, cut to a shorter length, then reheated in the furnace for 60 minutes before they were rolled again towards targeted gauge thickness. Total reduction of 96% was achieved for both alloys.
  • Hot rolled data represents properties of the sheets corresponding to the as-produced state in a case of Thin Slab Production including solidification, hot rolling, and coiling.
  • Cold rolling was applied to hot rolled sheet to reduce sheet thickness to 2 mm leading to significant strengthening of the sheet material through the Dynamic Nanophase Strengthening mechanism.
  • Subsequent heat treatment of the hot rolled and cold rolled sheet provides properties with strength of 1000 to 1200 MPa and ductility in the range from 17 to 24%. Final properties can vary depending on alloy chemistry as well as casting and post-processing parameters.
  • Structure #1a Homogenized Modal Structure transforms directly into Structure #2 (Nanomodal Structure) after a specific number of cycles of Mechanism #0 (Dynamic Nanophase Refinement) or if an additional heat treatment is needed to activate Mechanism #1 (Static Nanophase Refinement) to form Structure #2 (Nanomodal Structure).
  • Mechanism #0 Dynamic Nanophase Refinement
  • Mechanism #1 Static Nanophase Refinement
  • Mechanism #2 Dynamic Nanophase Strengthening
  • Plates were cast with different thicknesses in a range from 5 to 50 mm using an Indutherm VTC 800 V caster.
  • charges of different masses were weighed out for particular alloys according to the atomic ratios provided in Table 4.
  • the charges for Alloy 1 and Alloy 2 according to the atomic ratios provided in Table 4 were then placed into the crucible of an Indutherm VTC 800 V Tilt Vacuum Caster.
  • the feedstock was melted using RF induction and then poured into a copper die designed for casting plates with dimensions described in Table 13. All plates from each alloy were subjected to hot rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace.
  • the plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling. The plates were then repeatedly rolled down to 1.2 to 1.4 mm thickness. To mimic possible post-processing of the sheet produced by the Thin Slab Process, additional cold rolling with 39% reduction was applied to hot rolled plates with subsequent heat treatment at 1150° C. for 2 hrs.
  • the tensile specimens were cut from the rolled and heat treated sheets from Alloy 1 and Alloy 2 using a Brother HS-3100 wire electrical discharge machining (EDM).
  • the tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. Video extensometer was utilized for strain measurements. Tensile data for both alloys are plotted in FIG. 28 . Consistent properties with similar strength and ductility in the range from 20 to 29% for Alloy 1 and from 19 to 26% for Alloy 2 were measured in post-processed sheets independently from the as-cast thickness.
  • This Case Example demonstrates that Homogenized Modal Structure (Structure #1a, FIG. 8 ) forms in the Alloy 1 and Alloy 2 plates during hot rolling through Dynamic Nanophase Refinement (Mechanism #0, FIG. 8 ) resulting in the consistent properties independently from initial cast thickness. That is, provided one starts with Modal Structure, and undergoes Dynamic Nanophase Refinement to Homogenized Modal Structure, one can then continue with the sequence shown in FIG. 8 to achieve useful mechanical properties, regardless of the thickness of the initial cast thickness present in Structure 1 (i.e. when the thickness of the Modal Structure is greater than or equal to 2.0 mm, such as a thickness of greater than or equal to 2.0 mm to a thickness of 500 mm).
  • Plates with thicknesses of 20 mm were cast from Alloy 2 using an Indutherm VTC 800 V Tilt Vacuum Caster. Using commercial purity feedstock, charges of different masses were weighed out for particular alloy according to the atomic ratios provided in Table 4. The charges were then placed into the crucible of the caster. The feedstock was melted using RF induction and then poured into a copper die designed for casting plates with thickness of 20 mm. Cast plate was subjected to hot rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace. The plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling.
  • the plates were then hot rolled with multiple passes of 10% to 25% reduction mimicking multi-stand hot rolling during Stage 2 at the Thin Slab Process ( FIG. 2 ) or hot rolling process at Thick Slab Casting ( FIG. 1 ). Total hot rolling reduction was 88%. After hot rolling, the resultant sheet was heat treated at 950° C. for 6 hrs.
  • FIG. 29 shows the microstructure of the sheet after hot rolling with 88% reduction. It can be seen that hot rolling resulted in structural homogenization leading to formation of Homogenized Modal Structure (Structure #1a, FIG. 8 ) through Dynamic Nanophase Refinement (Mechanism #0, FIG. 8 ). However, while in the outer layer region, the fine boride phase is relatively uniform in size and homogeneously distributed in matrix, in the central layer region, although the boride phase is effectively broken up by the hot rolling, the distribution of boride phase is less homogeneous as at the outer layer. It can be seen that the boride distribution is not homogeneous. After an additional heat treatment at 950° C. for 6 hrs, as shown in FIG.
  • the boride phase is homogeneously distributed at both the outer layer and the central layer regions. In addition, the boride becomes more uniform in size. Comparison between FIG. 29 and FIG. 30 also suggests that the aspect ratio of the boride phase is smaller after heat treatment, its morphology is close to spherical geometry, and the boride size is more uniform through the sheet volume after heat treatment. The microstructure after the additional heat treatment is typical for the Nanomodal Structure (Structure #2, FIG. 8 ).
  • the heat treated sheet samples transform into the High Strength Nanomodal Structure during tensile testing resulting in an ultimate tensile strength (UTS) of 1222 MPa and a tensile elongation of 26.2% as compared to the UTS of 1193 MPa, and elongation of 17.9% before the heat treatment, underlining the effectiveness of the heat treatment on structural optimization.
  • UTS ultimate tensile strength
  • Tensile specimens were cut from the rolled and heat treated sheets from Alloy 8 using a Brother HS-3100 wire electrical discharge machining (EDM). The tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. Video extensometer was utilized for strain measurements. Tensile data for Alloy 8 after heat treatment at different conditions are plotted in FIG. 31 a . Tensile properties of Alloy 8 are shown to improve with additional hot rolling and heat treatment. Following 96% thickness reduction by hot rolling, the tensile elongation is >10% with tensile strength of approximately 1300 MPa.
  • Alloy 8 that has been heat treated at the HT3 condition possess tensile elongation of >15% with tensile strength approximately 1300 MPa.
  • FIG. 31 b illustrates the representative stress-strain curves showing alloy behavior improvement by increasing hot rolling reduction with subsequent heat treatment.
  • Tensile specimens were cut from the rolled and heat treated sheets from Alloy 16 using a Brother HS-3100 wire electrical discharge machining (EDM). The tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. Video extensometer was utilized for strain measurements. Tensile data for Alloy 16 after heat treatment at different conditions are plotted in FIG. 32 . Tensile properties of Alloy 16 are shown to improve with additional hot rolling and heat treatment.
  • Tensile specimens were cut from the rolled and heat treated sheets from Alloy 24 using a Brother HS-3100 wire electrical discharge machining (EDM). The tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. Video extensometer was utilized for strain measurements. Tensile data for Alloy 24 after heat treatment at different conditions are plotted in FIG. 33 a . Tensile properties of Alloy 24 are shown to improve with additional hot rolling and heat treatment.
  • FIG. 33 b illustrates the representative stress-strain curves showing alloy ductility improvement by increasing temperature of heat treatment after hot rolling with decreasing ductility.
  • TEM specimen preparation procedure includes cutting, thinning, electropolishing. First, samples were cut with electric discharge machine, and then thinned by grinding with pads of reduced grit size every time.
  • the TEM studies were done using a JEOL 2100 high-resolution microscope operated at 200 kV.
  • the TEM image of the microstructure in the Alloy 1 plate after hot rolling and heat treatment before deformation is shown in FIG. 34 . It can be seen that the Alloy 1 slab sample shows a textured microstructure due to hot rolling. Microstructure refinement is also seen in the sample. Since the sample was heat treated prior to the tensile deformation, the microstructure refinement indicates that Static Nanophase Refinement (Mechanism #1, FIG. 8 ) occurs during the heat treatment leading to Nanomodal Structure (Structure #2, FIG. 8 ) formation.
  • Mechanism #1, FIG. 8 Static Nanophase Refinement
  • the hot rolling prior heat treatment resulted in homogeneous distribution of the boride phase in matrix when Homogenized Modal Structure (Structure #1a, FIG. 8 ) was formed.
  • the Homogenized Modal Structure in this alloy corresponds to Type 2 (Table 3).
  • matrix grains of 200 to 500 nm in size can be found in the sample after heat treatment. Within the matrix grains, stacking faults can also be found, suggesting formation of austenite phase.
  • FIG. 35 shows the bright-field TEM images of the samples taken from the gage section of tensile specimens.
  • Mechanism #2, FIG. 8 Dynamic Nanophase Strengthening
  • Structure #3, FIG. 8 High Strength Nanomodal Structure
  • Grains of 200 to 300 nm in size are commonly observed in the matrix and very fine precipitates of hexagonal phases can be found.
  • the stacking faults shown in the samples before deformation disappeared after the tensile deformation, suggesting the austenite transforms to ferrite, and dislocations are generated in the matrix grains during the tensile deformation.
  • This Case Example illustrates High Strength Nanomodal Structure formation (Structure #3, FIG. 8 ) in Alloy 1 initially cast at 50 mm thickness with subsequent hot rolling and heat treatment. Structural development through enabling mechanisms follows the pathway illustrated in FIG. 8 .
  • TEM specimen preparation procedure includes cutting, thinning, electropolishing. First, samples were cut with electric discharge machine (EDM), and then thinned by grinding with pads of reduced grit size every time.
  • EDM electric discharge machine
  • the TEM image of the microstructure in the Alloy 8 plate after hot rolling and heat treatment before deformation is shown in FIG. 36 a .
  • the Alloy 8 sample before deformation shows a refined microstructure, as grains of several hundred nanometers are found in the sample confirming Homogenized Modal Structure (Structure 1a, FIG. 8 ) formation followed by Static Nanophase Refinement (Mechanism #1, FIG. 8 ) activation during heat treatment with formation of Nanomodal Structure (Structure #2, FIG. 8 ).
  • a modulation of dark and bright contrast is shown in the matrix grains, similar to the lamellar type structure.
  • FIG. 36 b shows slightly dark contrast showing incipient nano-size precipitates can be barely seen in the matrix prior to deformation. After deformation, the nano-size precipitates seem to develop a stronger contrast, as shown in FIG. 36 b .
  • the change of nano-size precipitates is better revealed by high magnification images.
  • FIG. 37 shows the matrix structure before and after deformation at a higher magnification. In contrast to the weak contrast shown by the nano-size precipitates before deformation, as it can be seen in FIG. 37 , the precipitates are better developed after deformation.
  • FIG. 37 b A close view of the precipitate regions suggests that they are composed of several smaller precipitates, FIG. 37 b .
  • Study by high-resolution TEM further reveals the structure of the nano-size precipitates.
  • FIG. 38 the lattice of nano-size precipitates is distinguished from the matrix, but their geometry is not clearly defined, suggesting that they might be just formed and perhaps in coherence with the matrix.
  • the precipitates are well identifiable with a size of generally 5 nm or less.
  • This Case Example illustrates High Strength Nanomodal Structure formation (Structure #3, FIG. 8 ) in Alloy 8 initially cast at 50 mm thickness with subsequent hot rolling and heat treatment. Structural development through the mechanisms follows the pathway illustrated in FIG. 8 .
  • TEM specimen preparation procedure includes cutting, thinning, electropolishing. First, samples were cut with electric discharge machine, and then thinned by grinding with pads of reduced grit size every time.
  • the TEM image of the Alloy 16 slab sample before deformation is shown in FIG. 39 a . It can be seen that the Alloy 16 slab sample shows a textured microstructure due to hot rolling. The rolling texture is further revealed by dark-field TEM image shown in FIG. 39 b . However, microstructure refinement is seen in the sample. As shown by both the bright-field and dark-field images, the refined grains of several hundred nanometers can be seen in the sample indicating that Static Nanophase Refinement (Mechanism #1, FIG. 8 ) occurs during the heat treatment leading to Nanomodal Structure (Structure #2, FIG. 8 ) formation. As shown in FIG. 39 b , matrix grains of 200 to 500 nm in size can be found in the sample after heat treatment.
  • Mechanism #1, FIG. 8 Static Nanophase Refinement
  • FIG. 39 b matrix grains of 200 to 500 nm in size can be found in the sample after heat treatment.
  • FIG. 40 shows the bright-field and dark-field TEM images of the samples made from the gage section of tensile specimen.
  • grains of 200 to 300 nm in size are commonly observed, and very fine precipitates of the new hexagonal phases can be found confirming that Dynamic Nanophase Strengthening (Mechanism #2) with formation of High Strength Nanomodal Structure (Structure #3) occurred during deformation. Additionally, dislocations are generated in the matrix grains during the tensile deformation.
  • This Case Example illustrates High Strength Nanomodal Structure formation (Structure #3, FIG. 8 ) in Alloy 16 initially cast at 50 mm thickness with subsequent hot rolling and heat treatment. Structural development through the mechanisms follows the pathway illustrated in FIG. 8 .
  • Plates with 50 mm thickness from Alloy 32 and Alloy 42 were cast using a Indutherm VTC 800 V Tilt Vacuum Caster was utilized to mimic the Stage 1 of the Thin Slab Process ( FIG. 2 ).
  • the plates from each alloy were subjected to hot rolling using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace.
  • the plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling.
  • the plates were then repeatedly rolled at between 10% and 25% reduction per pass down to 2 mm thickness mimicking multi-stand hot rolling at Stage 2 during the Thin Slab Process ( FIG. 2 ).
  • the plates were placed in the furnace for 1 to 2 min between rolling steps to allow then to return to temperature. If the plates became too long to fit in the furnace they were cooled, cut to a shorter length, then reheated in the furnace for 60 minutes before they were rolled again towards targeted gauge thickness. Total reduction at the hot rolling was 96%. Hot rolled sheets from both alloys were heat treated at 850° C. for 6 hr with slow cooling with furnace (0.75° C./min) to 500° C. with subsequent air cooling.
  • the tensile specimens were cut from the rolled and heat treated sheets from Alloy 32 and Alloy 42 using a Brother HS-3100 wire electrical discharge machining (EDM).
  • EDM wire electrical discharge machining
  • the tensile properties were tested 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 rigid and the top fixture moving with the load cell attached to the top fixture. A video extensometer was utilized for strain measurements.
  • Hot rolled data represents properties of the sheets corresponding to as-produced state in a case of Thin Slab Production including solidification, hot rolling and coiling (open symbols in FIG. 41 ). Both alloys show similar properties in hot rolled state with high ductility in the range from 45 to 48%. Heat treatment of the Alloy 42 sheet has changed the properties slightly while Alloy 32 has demonstrated a significant increase in ductility (up to 66.56%) in the heat treated state (solid symbols in FIG. 41 ) which may be due to elimination of defects and additional matrix grain coarsening.
  • Alloy 24 plate initially cast at 50 mm thickness was studied by TEM.
  • the casting was done using a Indutherm VTC 800 V Tilt Vacuum Caster, and then the slab was hot rolled to 2 mm thick sheet at 1100° C.
  • samples from Alloy 24 in the as-cast and hot rolled conditions were studied by TEM.
  • TEM specimen preparation procedure includes cutting, thinning, and electropolishing. First, samples were cut with electric discharge machine, and then thinned by grinding with pads of reduced grit size every time. Further thinning to 60 to 70 ⁇ m thickness was done by polishing with 9 ⁇ m, 3 ⁇ m and 1 ⁇ m diamond suspension solution respectively. Discs of 3 mm in diameter were punched from the foils and the final polishing was fulfilled with electropolishing using a twin-jet polisher. The chemical solution used was a 30% nitric acid mixed in a methanol base.
  • the TEM specimens were ion-milled using a Gatan Precision Ion Polishing System (PIPS). The ion-milling was done at 4.5 keV, and the inclination angle was reduced from 4° to 2° to open up the thin area.
  • the TEM studies were done using a JEOL 2100 high-resolution microscope operated at 200 kV.
  • FIG. 42 is the Modal Structure (Structure #1, FIG. 8 ).
  • the boride phase is long and slim, aligned at grain boundaries of matrix.
  • the size of boride phase can range from 1 ⁇ m to up to 10 ⁇ m, while the size of the matrix in between is typically 5 to 10 ⁇ m.
  • the boride phase resides at grain boundaries of matrix that fits the basic characteristic of the Modal Structure.
  • Partial transformation into the Nanomodal Structure (Structure #2, FIG. 8 ) in some areas can also be observed in this alloy as shown in FIG. 42 b where the matrix grains undergo refinement. Partial transformation might be related to slow cooling rate when alloy cast at large thicknesses resulting in extended time at elevated temperature to allow limited Static Nanophase Refinement (Mechanism #1, FIG. 8 ) in some areas.
  • the boride phase was broken up into small particles and is well scattered in the matrix indicating structural homogenization through Dynamic Nanophase Refinement (Mechanism #0, FIG. 8 ) leading to Homogenized Modal Structure formation (Structure #1a, FIG. 8 ).
  • the size of boride phase can be somewhere from 1 ⁇ m to 5 ⁇ m, but the slim geometry is largely reduced to a smaller aspect ratio.
  • the matrix grains, compared to the as-cast state, are significantly refined with the grain size of matrix reduced to 200 to 500 nm.
  • the matrix grains are elongated, aligning along the rolling direction after the rolling.
  • Elastic Modulus was measured for selected alloys listed in Table 22. Each alloy used was cast into a plate with thickness of 50 mm. Using a high temperature inert gas furnace the material was brought to the desired temperature, depending on alloy solidus temperature, prior to hot rolling. Initial hot rolling reduced the material thickness by approximately 85%. The oxide layer was removed from the hot rolled material using abrasive media. The center was sectioned from the resulting slab and hot rolled approximately an additional 75%. After removing the final oxide layer ASTM E8 subsize tensile samples were cut from center of the resulting material using wire electrical discharge machining (EDM). Tensile testing was performed on an Instron Model 3369 mechanical testing frame, using the Instron Bluehill control and analysis software.
  • EDM wire electrical discharge machining
  • 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. Tensile loading was performed to a load less than the yield point previously observed in tensile testing of the material, and this loading curve was used to obtain modulus values. Samples were pre-cycled under a tensile load below that of the predicted yield load to minimize the impact of grip settling on the measurements. Elastic modulus data in Table 23 is reported as an average value of 5 separate measurements. Modulus values vary in a range from 190 to 210 GPa typical for commercial steels and depend on alloy chemistry and thermo-mechanical treatment.
  • FIG. 44 shows the shrinkage funnel at the top of the figure.
  • the results of the chemical analysis are shown in FIG. 45 .
  • the content of each individual element in wt% is shown for the tested locations at the top (A) and bottom (B) of the cast plate for the four alloys identified.
  • the difference between the top (A) and bottom (B) ranges from 0.00 wt % to 0.19 wt % with no evidence for macrosegregation.
  • Tensile properties of selected alloys from Table 4 were compared with tensile properties of existing steel grades. The selected alloys and corresponding parameters are listed in Table 24. Tensile stress-strain curves are compared to that of existing Dual Phase (DP) steels ( FIG. 46 ); Complex Phase (CP) steels ( FIG. 47 ); Transformation Induced Plasticity (TRIP) steels ( FIG. 48 ); and Martensitic (MS) steels ( FIG. 49 ).
  • DP Dual Phase
  • CP Complex Phase
  • TRIP Transformation Induced Plasticity
  • MS Martensitic
  • a Dual Phase Steel may be understood as a steel type containing a ferritic matrix containing hard martensitic second phases in the form of islands
  • a Complex Phase Steel may be understood as a steel type containing a matrix consisting of ferrite and bainite containing small amounts of martensite, retained austenite, and pearlite
  • a Transformation Induced Plasticity steel may be understood as a steel type which consists of austenite embedded in a ferrite matrix which additionally contains hard bainitic and martensitic second phases
  • a Martensitic steel may be understood as a steel type consisting of a martensitic matrix which may contain small amounts of ferrite and/or bainite.
  • Example demonstrates that the alloys disclosed here have relatively superior mechanical properties as compared to existing advanced high strength (AHSS) steel grades with.
  • Ductility of 20% and above demonstrated by selected alloys provides cold formability of the sheet material and make it applicable to many processes such as for example cold stamping of a relatively complex part.
  • the plates were placed in a furnace pre-heated to 1140° C. for 60 minutes prior to the start of rolling.
  • the plates were then repeatedly rolled at between 10% and 25% reduction per pass down to 3.5 mm thickness mimicking multi-stand hot rolling at Stage 2 during the Thin Slab Process ( FIG. 2 ) or hot rolling step at Thick Slab Casting ( FIG. 1 ).
  • the plates were placed in the furnace for 1 to 2 min between rolling steps to allow them to return to temperature. If the plates became too long to fit in the furnace they were cooled, cut to a shorter length, then reheated in the furnace for 60 minutes before they were rolled again towards targeted gauge thickness. Total reduction of 96% was achieved for all alloys.
  • the chemical composition of the boron-free alloys herein (Alloy 63 through Alloy 74) is listed in Table 4 which provides the preferred atomic ratios utilized. These chemistries have been used for material processing through slab casting in an Indutherm VTC800V vacuum tilt casting machine. Alloys of designated compositions were weighed out in 3 kilogram charges using designated quantities of commercially-available ferroadditive powders of known composition and impurity content, and additional alloying elements as needed, according to the atomic ratios provided in Table 4 for each alloy. Weighed out Alloy charges were placed in zirconia coated silica-based crucibles and loaded into the casting machine. Melting took place under vacuum using a 14 kHz RF induction coil.
  • the 50 mm thick laboratory slab from each alloy was subjected to hot rolling at the temperature of 1250° C. except that from Alloy 68 which was rolled at 1250° C.
  • Rolling was done on a Fenn Model 061 single stage rolling mill, employing an in-line Lucifer EHS3GT-B18 tunnel furnace. Material was held at hot rolling temperature for an initial dwell time of 40 minutes to ensure homogeneous temperature. After each pass on the rolling mill, the sample was returned to the tunnel furnace with a 4 minute temperature recovery hold to correct for temperature lost during the hot rolling pass.
  • Hot rolling was conducted in two campaigns, with the first campaign achieving approximately 80% to 88% total reduction to a thickness of between 6 mm and 9.5 mm.
  • the density of the alloys was measured on-sections of cast material that had been hot rolled to between 6 mm and 9.5 mm. Sections were cut to 25 mm ⁇ 25 mm dimensions, and then surface ground to remove oxide from the hot rolling process. Measurements of bulk density were taken from these ground samples, 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 28 and was found to vary from 7.64 to 7.80 g/cm 3 . Experimental results have revealed that the accuracy of this technique is ⁇ 0.01 g/cm 3 .
  • the fully hot-rolled sheet was then subjected to cold rolling in multiple passes. Rolling was done on a Fenn Model 061 single stage rolling mill. A list of specific cold rolling parameters used for the alloys is shown in Table 29.
  • Tensile specimens were tested in the hot rolled, cold rolled, and heat treated conditions. 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 rigid and the top fixture moving; the load cell is attached to the top fixture.
  • Tensile properties of the alloys in the as hot rolled condition are listed in Table 31.
  • the ultimate tensile strength values may vary from 947 to 1329 MPa with tensile elongation from 20.5 to 55.4%.
  • the yield stress is in a range from 267 to 520 MPa.
  • the mechanical characteristic values in the steel alloys herein will depend on alloy chemistry and hot rolling conditions.
  • An example stress-strain curve for Alloy 63 in as hot rolled state is shown in FIG. 52 demonstrating typical Class 2 behavior ( FIG. 7 ).
  • Tensile properties of selected alloys after hot rolling and subsequent cold rolling are listed in Table 32 which represent Structure #3 or the High Strength Nanomodal Structure.
  • the ultimate tensile strength values may vary from 1402 to 1766 MPa with tensile elongation from 9.7 to 29.1%.
  • the yield stress is in a range from 913 to 1278 MPa.
  • the mechanical characteristic values in the steel alloys herein will depend on alloy chemistry and processing conditions.
  • Tensile properties of the hot rolled sheets after hot rolling with subsequent heat treatment at different parameters are listed in Table 33.
  • the ultimate tensile strength values may vary from 669 to 1352 MPa with tensile elongation from 15.9% to 78.1%.
  • the yield stress is in a range from 217 to 621 MPa.
  • the mechanical characteristic values in the steel alloys herein will depend on alloy chemistry and processing conditions.
  • Alloy 65 Plate with 50 mm thickness from Alloy 65 was cast in an Indutherm VTC800V vacuum tilt casting machine. Alloy of designated composition was weighed out in 3 kilogram charges using designated quantities of commercially-available ferroadditive powders of known composition and impurity content, and additional alloying elements as needed, according to the atomic ratios provided in Table 4. Weighed out Alloy charge was placed in zirconia coated silica-based crucibles and loaded into the casting machine. Melting took place under vacuum using a 14 kHz RF induction coil. Alloy charge was heated until fully molten, with a period of time between 45 seconds and 60 seconds after the last point at which solid constituents were observed, in order to provide superheat and ensure melt homogeneity. Melt was then poured into a water-cooled copper die to form laboratory cast slab of approximately 50 mm thick which is in the thickness range for the Thin Slab Casting process and 75 mm ⁇ 100 mm in size.
  • the 50 mm thick laboratory slab from the Alloy 65 was subjected to hot rolling at the temperature of 1250° C. with a total reduction of 97%.
  • the fully hot-rolled sheet was then subjected to cold rolling in multiple passes down to thickness of 1.2 mm.
  • Cold rolled sheet was heat treated at 850° C. for 5 minutes that mimic in-line annealing at commercial sheet production.
  • SEM specimens the cross-sections of the sheet sample in as-cast state, after hot rolling, and after cold rolling with subsequent heat treatment were cut and ground by SiC paper and then polished progressively with diamond media paste down to 1 ⁇ m grit. The final polishing was done with 0.02 ⁇ m grit SiO 2 solution.
  • Microstructures of samples from Alloy 65 were examined by scanning electron microscopy (SEM) using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.
  • FIG. 53 shows SEM images of microstructure in Alloy 65 in as-cast state, after hot rolling, and after cold rolling with subsequent heat treatment demonstrating a structural development from Modal Structure in as-cast state ( FIG. 53 a ), Nanomodal Structure in the hot rolled state ( FIG. 53 b ), and High Strength Nanomodal Structure after cold rolling ( FIG. 53 c ).

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US9493855B2 (en) * 2013-02-22 2016-11-15 The Nanosteel Company, Inc. Class of warm forming advanced high strength steel
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US20150090372A1 (en) * 2013-10-02 2015-04-02 The Nanosteel Company, Inc. Recrystallization, Refinement, and Strengthening Mechanisms For Production Of Advanced High Strength Metal Alloys
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US11560605B2 (en) 2019-02-13 2023-01-24 United States Steel Corporation High yield strength steel with mechanical properties maintained or enhanced via thermal treatment optionally provided during galvanization coating operations
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