WO2018160387A1 - Formabilité de bord améliorée dans les alliages métalliques - Google Patents

Formabilité de bord améliorée dans les alliages métalliques Download PDF

Info

Publication number
WO2018160387A1
WO2018160387A1 PCT/US2018/018751 US2018018751W WO2018160387A1 WO 2018160387 A1 WO2018160387 A1 WO 2018160387A1 US 2018018751 W US2018018751 W US 2018018751W WO 2018160387 A1 WO2018160387 A1 WO 2018160387A1
Authority
WO
WIPO (PCT)
Prior art keywords
alloy
edge
hole
tensile
punched
Prior art date
Application number
PCT/US2018/018751
Other languages
English (en)
Inventor
Daniel James Branagan
Andrew E. Frerichs
Brian E. Meacham
Grant G. Justice
Andrew T. Ball
Jason K. Walleser
Kurtis R. CLARK
Logan J. TEW
Scott T. ANDERSON
Scott T. LARISH
Sheng Cheng
Taylor L. Giddens
Alla V. Sergueeva
Original Assignee
The Nanosteel Company, Inc.
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Priority claimed from US15/438,313 external-priority patent/US10465260B2/en
Application filed by The Nanosteel Company, Inc. filed Critical The Nanosteel Company, Inc.
Priority to CA3053383A priority Critical patent/CA3053383A1/fr
Priority to EP18760610.8A priority patent/EP3585532A4/fr
Priority to JP2019544895A priority patent/JP2020509233A/ja
Priority to CN201880013756.6A priority patent/CN110382130A/zh
Priority to KR1020197027312A priority patent/KR20190130131A/ko
Publication of WO2018160387A1 publication Critical patent/WO2018160387A1/fr
Priority to JP2023039718A priority patent/JP2023075277A/ja

Links

Classifications

    • 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
    • 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
    • C21D1/26—Methods of annealing
    • 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
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • 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
    • C22C38/004—Very low carbon steels, i.e. having a carbon content of less than 0,01%
    • 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
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium 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
    • C22C38/32—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
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • 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
    • 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
    • C21D2211/001—Austenite
    • 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • This disclosure relates to methods for mechanical property improvement in a metallic alloy that has undergone one or more mechanical property losses as a consequence of shearing, such as in the formation of a sheared edge portion or a punched hole. More specifically, methods are disclosed that provide the ability to improve mechanical properties of metallic alloys that have been formed with one or more sheared edges which may otherwise serve as a limiting factor for industrial applications.
  • LSS Low Strength Steels
  • HSS High Strength Steels
  • AHSS Advanced High Strength Steels
  • High-Strength Steels are classified as exhibiting ultimate tensile strengths from 270 to 700 MPa and include such types as high strength low alloy, high strength interstitial free and bake hardenable steels.
  • Advanced High-Strength Steels (AHSS) steels are classified by ultimate tensile strengths greater than 700 MPa and include such types as Martensitic steels (MS), Dual Phase (DP) steels, Transformation Induced Plasticity (TRIP) steels, and Complex Phase (CP) steels.
  • MS Martensitic steels
  • DP Dual Phase
  • TRIP Transformation Induced Plasticity
  • CP Complex Phase
  • ductility tensile elongation
  • tensile elongation of LSS, HSS and AHSS ranges from 25% to 55%, 10% to 45%, and 4% to 30%, respectively.
  • AHSS advanced high strength steels
  • the high strength of AHSS allows for a designer to reduce the thickness of a finished part while still maintaining comparable or improved mechanical properties. In reducing the thickness of a part, less mass is needed to attain the same or better mechanical properties for the vehicle thereby improving vehicle fuel efficiency. This allows the designer to improve the fuel efficiency of a vehicle while not compromising on safety.
  • Formability is the ability of a material to be made into a particular geometry without cracking, rupturing or otherwise undergoing failure.
  • High formability steel provides benefit to a part designer by allowing for the creation of more complex part geometries allowing for reduction in weight.
  • Formability may be further broken into two distinct forms: edge formability and bulk formability.
  • Edge formability is the ability for an edge to be formed into a certain shape. Edges on materials are created through a variety of methods in industrial processes, including but not limited to punching, shearing, piercing, stamping, perforating, cutting, or cropping.
  • the devices used to create these edges are as diverse as the methods, including but not limited to various types of mechanical presses, hydraulic presses, and/or electromagnetic presses.
  • the range of speeds for edge creation is also widely varying, with speeds as low as 0.25 mm/s and as high as 3700 mm s.
  • the wide variety of edge forming methods, devices, and speeds results in a myriad of different edge conditions in use commercially today.
  • Edges being free surfaces, are dominated by defects such as cracks or structural changes in the sheet resulting from the creation of the sheet edge. These defects adversely affect the edge formability during forming operations, leading to a decrease in effective ductility at the edge.
  • Bulk formability on the other hand is dominated by the intrinsic ductility, structure, and associated stress state of the metal during the forming operation. Bulk formability is affected primarily by available deformation mechanisms such as dislocations, twinning, and phase transformations. Bulk formability is maximized when these available deformation mechanisms are saturated within the material, with improved bulk formability resulting from an increased number and availability of these mechanisms.
  • Edge formability can be measured through hole expansion measurements, whereby a hole is made in a sheet and that hole is expanded by means of a conical punch.
  • Previous studies have shown that conventional AHSS materials suffer from reduced edge formability compared with other LSS and HSS when measured by hole expansion [M.S. Billur, T. Altan, "Challenges in forming advanced high strength steels", Proceedings of New Developments in Sheet Metal Forming, pp.285-304, 2012].
  • Dual Phase (DP) steels with ultimate tensile strength of 780 MPa achieve less than 20% hole expansion
  • Interstitial Free steels (IF) with ultimate tensile strength of approximately 400 MPa achieve around 100% hole expansion ratio.
  • This reduced edge formability complicates adoption of AHSS in automotive applications, despite possessing desirable bulk formability.
  • a method for improving one or more mechanical properties in a metallic alloy that has undergone a mechanical property loss as a consequence of the formation of one or more sheared edges comprising:
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said alloy and cooling at a rate of ⁇ 250 K/s or solidifying to a thickness of > 2.0 mm up to 500 mm and forming an alloy having a T m and matrix grains of 2 ⁇ to 10,000 ⁇ ;
  • the present disclosure also relates to a method for improving the hole expansion ratio in a metallic alloy that had undergone a hole expansion ratio loss as a consequence of forming a hole with a sheared edge comprising:
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said alloy and cooling at a rate of ⁇ 250 K/s or solidifying to a thickness of > 2.0 mm up to 500 mm and forming an alloy having a T m and matrix grains of 2 ⁇ to 10,000 ⁇ ;
  • the present invention also relates to method for improving the hole expansion ratio in a metallic alloy that had undergone a hole expansion ratio loss as a consequence of forming a hole with a sheared edge comprising:
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said alloy and cooling at a rate of ⁇ 250 K/s or solidifying to a thickness of > 2.0 mm up to 500 mm and forming an alloy having a Tm and matrix grains of 2 ⁇ to 10,000 ⁇ ; b. heating said alloy to a temperature of > 700 °C and below the Tm of said alloy and at a strain rate of 10 "6 to 10 4 and reducing said thickness of said alloy and providing a first resulting alloy having an ultimate tensile strength of 921 MPa to 1413 MPa and an elongation of 12.0% to 77.7%;
  • HER2 (0.01 to 0.30XHER ;
  • the present invention also relates to a method for punching one or more holes in a metallic alloy comprising:
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said alloy and cooling at a rate of ⁇ 250 K/s or solidifying to a thickness of > 2.0 mm up to 500 mm and forming an alloy having a Tm and matrix grains of 2 ⁇ to 10,000 ⁇ ; b. heating said alloy to a temperature of > 700 °C and below the Tm of said alloy and at a strain rate of 10 "6 to 10 4 and reducing said thickness of said alloy and providing a first resulting alloy having an ultimate tensile strength of 921 MPa to 1413 MPa and an elongation of 12.0% to 77.7%;
  • the present invention also relates to a method for expanding an edge in an alloy
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said alloy and cooling at a rate of ⁇ 250 K/s or solidifying to a thickness of > 2.0 mm up to 500 mm and forming an alloy having a Tm;
  • the present invention also relates to a method for expanding the edge of an alloy comprising:
  • a metal alloy comprising at least 50 atomic % iron and at least four or more elements selected from Si, Mn, B, Cr, Ni, Cu or C, wherein said alloy has an ultimate tensile strength of 799 MPa to 1683 MPa and an elongation of 6.6 to
  • FIG. 1A Structural pathway for the formation of High Strength Nanomodal Structure and associated mechanisms.
  • FIG. IB Structural pathway for the formation of Recrystallized Modal Structure
  • FIG. 2 Structural pathway toward developing Refined High Strength Nanomodal
  • FIG. 3 Images of laboratory cast 50 mm slabs from: a) Alloy 9 and b) Alloy 12.
  • FIG. 4 Images of hot rolled sheet after laboratory casting from: a) Alloy 9 and b) Alloy
  • FIG. 5 Images of cold rolled sheet after laboratory casting and hot rolling from: a) Alloy
  • FIG. 6 Microstructure of solidified Alloy 1 cast at 50 mm thickness: a) Backscattered
  • FIG. 7 X-ray diffraction pattern for the Modal Structure in Alloy 1 alloy after solidification: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 8 Microstructure of Alloy 1 after hot rolling to 1.7 mm thickness: a) Backscattered
  • FIG. 9 X-ray diffraction pattern for the Nanomodal Structure in Alloy 1 after hot rolling:
  • FIG. 10 Microstructure of Alloy 1 after cold rolling to 1.2 mm thickness: a) Backscattered
  • FIG. 11 X-ray diffraction pattern for the High Strength Nanomodal Structure in Alloy 1 after cold rolling: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 12 Bright-field TEM micrographs of microstructure in Alloy 1 after hot rolling, cold rolling and annealing at 850°C for 5 min exhibiting the Recrystallized Modal Structure: a) Low magnification image, b) High magnification image with selected electron diffraction pattern showing crystal structure of austenite phase.
  • FIG. 12 Bright-field TEM micrographs of microstructure in Alloy 1 after hot rolling, cold rolling and annealing at 850°C for 5 min exhibiting the Recrystallized Modal Structure: a) Low magnification image, b) High magnification image with selected electron diffraction pattern showing crystal structure of austenite phase.
  • FIG. 14 X-ray diffraction pattern for the Recrystallized Modal Structure in Alloy 1 after annealing: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 15 Bright-field TEM micrographs of microstructure in Alloy 1 showing Refined
  • FIG. 16 Backscattered SEM micrographs of microstructure in Alloy 1 showing Refined
  • High Strength Nanomodal Structure (Mixed Microconstituent Structure): a) Low magnification image, b) High magnification image.
  • FIG. 17 X-ray diffraction pattern for Refined High Strength Nanomodal Structure in
  • FIG. 18 Microstructure of solidified Alloy 2 cast at 50 mm thickness: a) Backscattered
  • FIG. 19 X-ray diffraction pattern for the Modal Structure in Alloy 2 after solidification: a)
  • FIG. 20 Microstructure of Alloy 2 after hot rolling to 1.7 mm thickness: a) Backscattered
  • FIG. 21 X-ray diffraction pattern for the Nanomodal Structure in Alloy 2 after hot rolling: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 22 Microstructure of Alloy 2 after cold rolling to 1.2 mm thickness: a) Backscattered
  • FIG. 23 X-ray diffraction pattern for the High Strength Nanomodal Structure in Alloy 2 after cold rolling: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 24 Bright-field TEM micrographs of microstructure in Alloy 2 after hot rolling, cold rolling and annealing at 850°C for 10 min exhibiting the Recrystallized Modal Structure: a) Low magnification image, b) High magnification image with selected electron diffraction pattern showing crystal structure of austenite phase.
  • FIG. 25 Backscattered SEM micrographs of microstructure in Alloy 2 after hot rolling, cold rolling and annealing at 850°C for 10 min exhibiting the Recrystallized Modal Structure: a) Low magnification image, b) High magnification image.
  • FIG. 26 X-ray diffraction pattern for the Recrystallized Modal Structure in Alloy 2 after annealing: a) Experimental data, b) Rietveld refinement analysis.
  • FIG. 27 Microstructure in Alloy 2 showing Refined High Strength Nanomodal Structure
  • FIG. 28 X-ray diffraction pattern for Refined High Strength Nanomodal Structure in
  • FIG. 29 Tensile properties of Alloy 1 at various stages of laboratory processing.
  • FIG. 30 Tensile results for Alloy 13 at various stages of laboratory processing.
  • FIG. 31 Tensile results for Alloy 17 at various stages of laboratory processing.
  • FIG. 32 Tensile properties of the sheet in hot rolled state and after each step of cold rolling/annealing cycles demonstrating full property reversibility at each cycle in: a) Alloy, b) Alloy 2.
  • FIG. 33 A bend test schematic showing a bending device with two supports and a former (International Organization for Standardization, 2005).
  • FIG. 34 Images of bend testing samples from Alloy 1 tested to 180°: a) Picture of a full set of samples tested to 180° without cracking, and b) A close-up view of the bend of a tested sample.
  • FIG. 35 a) Tensile test results of the punched and EDM cut specimens from selected alloys demonstrating property decrease due to punched edge damage, b) Tensile curves of the selected alloys for EDM cut specimens.
  • FIG. 36 SEM images of the specimen edges in Alloy 1 after a) EDM cutting and b)
  • FIG. 37 SEM images of the microstructure near the edge in Alloy 1: a) EDM cut specimens and b) Punched specimens.
  • FIG. 38 Tensile test results for punched specimens from Alloy 1 before and after annealing demonstrating full property recovery from edge damage by annealing. Data for EDM cut specimens for the same alloy are shown for reference.
  • FIG. 39 Example tensile stress-strain curves for punched specimens from Alloy 1 with and without annealing.
  • FIG. 40 Tensile stress-strain curves illustrating the response of cold rolled Alloy 1 to recovery temperatures in the range between 400°C and 850°C; a) Tensile curves, b) Yield strength.
  • FIG. 41 Bright-field TEM images of cold rolled ALLOY 1 samples exhibiting the highly deformed and textured High Strength Nanomodal Structure: a) Lower magnification image, b) Higher magnification image.
  • FIG. 42 Bright-field TEM images of ALLOY 1 samples annealed at 450°C 10 min exhibiting the highly deformed and textured High Strength Nanomodal Structure with no recrystallization occurred: a) Lower magnification image, b) Higher magnification image.
  • FIG. 43 Bright-field TEM images of ALLOY 1 samples annealed at 600°C 10 min exhibiting nanoscale grains signaling the beginning of recrystallization: a) Lower magnification image, b) Higher magnification image.
  • FIG. 44 Bright-field TEM images of ALLOY 1 samples annealed at 650°C 10 min exhibiting larger grains indicating the higher extent of recrystallization: a) Lower magnification image, b) Higher magnification image.
  • FIG. 45 Bright-field TEM images of ALLOY 1 samples annealed at 700°C 10 min exhibiting recrystallized grains with a small fraction of untransformed area, and electron diffraction shows the recrystallized grains are austenite: a) Lower magnification image, b) Higher magnification image.
  • FIG. 46 Model Time Temperature Transformation Diagram representing response of the steel alloys herein to temperature at annealing.
  • recovery mechanisms are activated.
  • recovery and recrystallization mechanisms are activated.
  • FIG. 47 Tensile properties of punched specimens before and after annealing at different temperatures: a) Alloyl, b) Alloy 9, and c) Alloy 12.
  • FIG. 48 Schematic illustration of the sample position for structural analysis.
  • FIG. 49 Alloy 1 punched E8 samples in the as-punched condition: a) Low magnification image showing a triangular deformation zone at the punched edge which is located on the right side of the picture. Additionally close up areas for the subsequent micrographs are provided, b) Higher magnification image showing the deformation zone, c) Higher magnification image showing the recrystallized structure far away from the deformation zone, d) Higher magnification image showing the deformed structure in the deformation zone.
  • FIG. 50 Alloy 1 punched E8 samples after annealing at 650°C for 10 min: a) Low magnification image showing the deformation zone at edge, punching in upright direction. Additionally, close up areas for the subsequent micrographs are provided: b) Higher magnification image showing the deformation zone, c) Higher magnification image showing the recrystallized structure far away from the deformation zone, d) Higher magnification image showing the recovered structure in the deformation zone.
  • FIG. 51 Alloy 1 punched E8 samples after annealing at 700°C for 10 min: a) Low magnification image showing the deformation zone at edge, punching in upright direction. Additionally, close up areas for the subsequent micrographs are provided, b) Higher magnification image showing the deformation zone, c) Higher magnification image showing the recrystallized structure far away from the deformation zone, d) Higher magnification image showing the recrystallized structure in the deformation zone.
  • FIG. 52 Tensile properties for specimens punched at varied speeds from: a) Alloy 1, b)
  • FIG. 53 HER results for Alloy 1 in a case of punched vs milled hole.
  • FIG. 54 Cutting plan for SEM microscopy and microhardness measurement samples from
  • FIG. 55 A schematic illustration of microhardness measurement locations.
  • FIG. 56 Microhardness measurement profile in Alloy 1 HER tested samples with: a)
  • FIG. 57 Microhardness profiles for Alloy 1 in various stages of processing and forming, demonstrating the progression of edge structure transformation during hole punching and expansion.
  • FIG. 58 Microhardness data for HER tested samples from Alloy 1 with punched and milled holes. Circles indicate a position of the TEM samples in respect to hole edge.
  • FIG. 59 Bright field TEM image of the microstructure in the Alloy 1 sheet sample before
  • FIG. 60 Bright field TEM micrographs of microstructure in the HER test sample from
  • FIG. 62 Focused Ion Beam (FIB) technique used for precise sampling near the edge of the punched hole in the Alloy 1 sample: a) FIB technique showing the general sample location of the milled TEM sample, b) Close up view of the cut-out TEM sample with indicated location from the hole edge.
  • FIB Focused Ion Beam
  • FIG. 63 Bright field TEM micrographs of microstructure in the sample from Alloy 1 with a punched hole at a location of -10 micron from the hole edge.
  • FIG. 64 Hole expansion ratio measurements for Alloy 1 with and without annealing of punched holes.
  • FIG. 65 Hole expansion ratio measurements for Alloy 9 with and without annealing of punched holes.
  • FIG. 66 Hole expansion ratio measurements for Alloy 12 with and without annealing of punched holes.
  • FIG. 67 Hole expansion ratio measurements for Alloy 13 with and without annealing of punched holes.
  • FIG. 68 Hole expansion ratio measurements for Alloy 17 with and without annealing of punched holes.
  • FIG. 69 Tensile performance of Alloy 1 tested with different edge conditions. Note that tensile samples with Punched edge condition have reduced tensile performance when compared to tensile samples with wire EDM cut and punched with subsequent annealing (850°C for 10 minutes) edge conditions.
  • FIG. 70 Edge formability as measured by hole expansion ratio response of Alloy 1 as a function of edge condition. Note that holes in the Punched condition have lower edge formability than holes in the wire EDM cut and punched with subsequent annealing (850°C for 10 minutes) conditions.
  • FIG. 71 Punch speed dependence of Alloy 1 edge formability as a function of punch speed, measured by hole expansion ratio. Note the consistent increase in hole expansion ratio with increasing punch speed.
  • FIG. 72 Punch speed dependence of Alloy 9 edge formability as a function of punch speed, measured by hole expansion ratio. Note the rapid increase in hole expansion ratio up to approximately 25 mm/s punch speed followed by a gradual increase in hole expansion ratio.
  • FIG. 73 Punch speed dependence of Alloy 12 edge formability as a function of punch speed, measured by hole expansion ratio. Note the rapid increase in hole expansion ratio up to approximately 25 mm/s punch speed followed by a continued increase in hole expansion ratio with punch speeds of >100 mm/s.
  • FIG. 74 Punch speed dependence of commercial Dual Phase 980 steel edge formability measured by hole expansion ratio. Note the hole expansion ratio is consistently 21% with + 3% variance for commercial Dual Phase 980 steel at all punch speeds tested.
  • FIG. 75 Schematic drawings of non-flat punch geometries: 6° taper (left), 7° conical
  • FIG. 76 Punch geometry effect on Alloy 1 at 28 mm/s, 114 mm/s, and 228 mm/s punch speed. Note that for the Alloy 1, the effect of punch geometry diminishes at 228 mm/s punch speed.
  • FIG. 77 Punch geometry effect on Alloy 9 at 28 mm/s, 114 mm/s, and 228 mm/s punch speeds. Note that the 7° conical punch and the conical flat punch result in the highest hole expansion ratio.
  • FIG. 78 Punch geometry effect on Alloy 12 at 28 mm/s, 114 mm/s, and 228 mm/s punch speed. Note that the 7° conical punch results at 228 mm/s punch speed in the highest hole expansion ratio measured for all alloys.
  • FIG. 79 Punch geometry effect on Alloy 1 at 228 mm/s punch speed. Note that all punch geometries result in nearly equal hole expansion ratios of approximately 21%.
  • FIG. 80 Hole punch speed dependence of commercial steel grades edge formability measured by hole expansion ratio.
  • FIG. 81 The post uniform elongation and hole expansion ratio correlation as predicted by [Paul S.K., J Mater Eng Perform 2014; 23:3610.] with data for selected commercial steel grades from the same paper along with Alloy 1 and Alloy 9 data.
  • FIG. 82 The measured hole expansion ratio in samples from Alloy 1 as a function of hole expansion speed.
  • FIG. 83 The measured hole expansion ratio in samples from Alloy 9 as a function of hole expansion speed.
  • FIG. 84 The measured hole expansion ratio in samples from Alloy 12 as a function of hole expansion speed.
  • FIG. 85 Images of the microstructure in the sheet from Alloy 9; a) SEM image of the microstructure, b) Higher magnification SEM image of the microstructure, c) Optical image of the etched surface, and d) Higher magnification optical image of the etched surface.
  • FIG. 86 The measured hole expansion ratio as a function of hole punching speed and hole expansion speed for sheet of Alloy 9.
  • FIG. 87 The average magnetic phases volume percent (Fe%) in the HER tested samples with different hole punching speed and hole expansion speed as a function of the distance from the hole edge.
  • FIG. 88 The measured hole expansion ratio in samples from Alloy 1, Alloy 9, and Alloy
  • FIG. 89 SEM images at low magnification of the cross section near the hole edge in the
  • FIG. 91 SEM images at low magnification of the cross section near the hole edge in the Alloy 1 samples with holes prepared by different methods after expansion during HER testing; a) Punched hole, b) EDM cut hole, c) Milled hole, and d) Laser cut hole.
  • FIG. 92 SEM images of sample cross sections near the hole edge after HER testing (i.e.
  • the steel alloys herein undergo a unique pathway of structural formation through specific mechanisms as illustrated in FIG. 1A and FIG. IB.
  • Initial structure formation begins with melting the alloy and cooling and solidifying and forming an alloy with Modal Structure (Structure #1, FIG. 1A).
  • the Modal Structure exhibits a primarily austenitic matrix (gamma-Fe) which may contain, depending on the specific alloy chemistry, ferrite grains (alpha-Fe), martensite, and precipitates including borides (if boron is present) and/or carbides (if carbon is present).
  • the grain size of the Modal Structure will depend on alloy chemistry and the solidification conditions. For example, thicker as-cast structures (e.g.
  • the Modal Structure preferably exhibits an austenitic matrix (gamma-Fe) with grain size and/or dendrite length from 2 to 10,000 ⁇ and precipitates at a size of 0.01 to 5.0 ⁇ in laboratory casting. Matrix grain size and precipitate size might be larger, up to a factor of 10 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters.
  • austenitic matrix gamma-Fe
  • Matrix grain size and precipitate size might be larger, up to a factor of 10 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters.
  • Steel alloys herein with the Modal Structure depending on starting thickness size and the specific alloy chemistry typically exhibits the following tensile properties, yield strength from 144 to 514 MPa, ultimate tensile strength in a range from 411 to 907 MPa, and total ductility from 3.7 to 24.4%.
  • Steel alloys herein with the Modal Structure can be homogenized and refined through the Nanophase Refinement (Mechanism #1, FIG. 1A) by exposing the steel alloy to one or more cycles of heat and stress ultimately leading to formation of the Nanomodal Structure (Structure #2, FIG. 1A).
  • the Modal Structure when formed at thickness of greater than or equal to 2.0 mm, or formed at a cooling rate of less than or equal to 250 K/s, is preferably heated to a temperature of 700°C to a temperature below the solidus temperature (T m ) and at strain rates of 10 "6 to 10 4 with a thickness reduction. Transformation to Structure #2 occurs in a continuous fashion through the intermediate Homogenized Modal Structure (Structure #la, FIG. 1A) as the steel alloy undergoes mechanical deformation during successive application of temperature and stress and thickness reduction such as what can be configured to occur during hot rolling.
  • the Nanomodal Structure (Structure #2, FIG. 1A) has a primary austenitic matrix (gamma- Fe) and, depending on chemistry, may additionally contain ferrite grains (alpha-Fe) and/or precipitates such as borides (if boron is present) and/or carbides (if carbon is present).
  • the Nanomodal Structure typically exhibits a primary austenitic matrix (gamma-Fe) with grain size of 1.0 to 100 ⁇ and/or precipitates at a size 1.0 to 200 nm in laboratory casting. Matrix grain size and precipitate size might be larger up to a factor of 5 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters.
  • Structure #2 is preferably formed at thickness of 1 mm to 500 mm.
  • the Dynamic Nanophase Strengthening Mechanism (Mechanism #2, FIG. 1A) is activated leading to formation of the High Strength Nanomodal Structure (Structure #3, FIG. 1A).
  • the stress is at a level above the alloy's respective yield strength in a range from 250 to 600 MPa depending on alloy chemistry.
  • the High Strength Nanomodal structure typically exhibits a ferritic matrix (alpha-Fe) which, depending on alloy chemistry, may additionally contain austenite grains (gamma-Fe) and precipitate grains which may include borides (if boron is present) and/or carbides (if carbon is present).
  • austenite grains gamma-Fe
  • precipitate grains which may include borides (if boron is present) and/or carbides (if carbon is present).
  • borides if boron is present
  • carbides if carbon is present
  • the strengthening transformation occurs during strain under applied stress that defines Mechanism #2 as a dynamic process during which the metastable austenitic phase (gamma-Fe) transforms into ferrite (alpha-Fe) with precipitates.
  • a fraction of the austenite will be stable and will not transform. Typically, as low as 5 volume percent and as high as 95 volume percent of the matrix will transform.
  • the High Strength Nanomodal Structure typically exhibits a ferritic matrix (alpha-Fe) with matrix grain size of 25 nm to 50 ⁇ and precipitate grains at a size of 1.0 to 200 nm in laboratory casting. Matrix grain size and precipitate size might be larger up to a factor of 2 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters. Steel alloys herein with the High Strength Nanomodal Structure typically exhibits the following tensile properties, yield strength from 718 to 1645 MPa, ultimate tensile strength in a range from 1356 to 1831 MPa, and total ductility from 1.6 to 32.8%. Structure #3 is preferably formed at thickness of 0.2 to 25.0 mm.
  • the High Strength Nanomodal Structure (Structure #3, FIG. 1A and FIG. IB) has a capability to undergo Recrystallization (Mechanism #3, FIG. IB) when subjected to heating below the melting point of the alloy with transformation of ferrite grains back into austenite leading to formation of Recrystallized Modal Structure (Structure #4, FIG. IB). Partial dissolution of nanoscale precipitates also takes place. Presence of borides and/or carbides is possible in the material depending on alloy chemistry. Preferred temperature ranges for a complete transformation occur from 650°C up to the T m of the specific alloy. When recrystallized, the Structure #4 contains few dislocations or twins and stacking faults can be found in some recrystallized grains.
  • the Recrystallized Modal Structure typically exhibits a primary austenitic matrix (gamma-Fe) with grain size of 0.5 to 50 ⁇ and precipitate grains at a size of 1.0 to 200 nm in laboratory casting. Matrix grain size and precipitate size might be larger up to a factor of 2 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters.
  • Steel alloys herein with the Recrystallized Modal Structure typically exhibit the following tensile properties: yield strength from 197 to 1372 MPa, ultimate tensile strength in a range from 799 to 1683 MPa, and total ductility from 10.6 to 86.7%.
  • the presence of borides (if boron is present) and/or carbides (if carbon is present) is possible in the material depending on alloy chemistry.
  • the untransformed part of the microstructure is represented by austenitic grains (gamma-Fe) with a size from 0.5 to 50 ⁇ and additionally may contain distributed precipitates with size of 1 to 200 nm. These highly deformed austenitic grains contain a relatively large number of dislocations due to existing dislocation processes occurring during deformation resulting in high fraction of dislocations (10 8 to 1010 mm - " 2 ).
  • the transformed part of the microstructure during deformation is represented by refined ferrite grains (alpha-Fe) with additional precipitate through Nanophase Refinement & Strengthening (Mechanism #4, FIG. IB).
  • the size of refined grains of ferrite varies from 50 to 2000 nm and size of precipitates is in a range from 1 to 200 nm in laboratory casting. Matrix grain size and precipitate size might be larger up to a factor of 2 at commercial production depending on alloy chemistry, starting casting thickness and specific processing parameters.
  • the size of the "pockets" of transformed and highly refined microstructure typically varies from 0.5 to 20 ⁇ .
  • the volume fraction of the transformed vs untransformed areas in the microstructure can be varied by changing the alloy chemistry including austenite stability from typically a 95:5 ratio to 5:95, respectively.
  • Steel alloys herein with the Refined High Strength Nanomodal Structure typically exhibit the following tensile properties: yield strength from 718 to 1645 MPa, ultimate tensile strength in a range from 1356 to 1831 MPa, and total ductility from 1.6 to 32.8%.
  • Final thicknesses of the material may therefore fall in the range from 0.2 to 25 mm.
  • cubic precipitates may be present in the steel alloys herein at all stages with a Fm3m (#225) space group.
  • Additional nanoscale precipitates may be formed as a result of deformation through Dynamic Nanophase Strengthening Mechanism (Mechanism #2) and/or Nanophase Refinement & Strengthening (Mechanism #4) that are represented by a dihexagonal pyramidal class hexagonal phase with a P6 3 m c space group (#186) and/or a ditrigonal dipyramidal class with a hexagonal P6bar2C space group (#190).
  • the precipitate nature and volume fraction depends on the alloy composition and processing history.
  • the size of nanoprecipitates can range from 1 nm to tens of nanometers, but in most cases below 20 nm. Volume fraction of precipitates is generally less than 20%.
  • Modal Structure (Structure #1) in steel alloys herein occurs during alloy solidification.
  • 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 lxlO 3 to lxlO 3 K/s.
  • the as-cast thickness will be dependent on the production method with Thin Slab Casting typically in the range of 20 to 150 mm in thickness and Thick Slab Casting typically in the range of 150 to 500 mm in thickness. Accordingly, as cast thickness may fall in the range of 20 to 500 mm, and at all values therein, in 1 mm increments. Accordingly, as cast thickness may be 21 mm, 22 mm, 23 mm, etc., up to 500 mm.
  • Hot rolling of solidified slabs from the alloys is the next processing step with production either of transfer bars in the case of Thick Slab Casting or coils in the case of Thin Slab Casting.
  • the Modal Structure transforms in a continuous fashion into a partial and then fully Homogenized Modal Structure (Structure #la) through Nanophase Refinement (Mechanism #1).
  • the Nanomodal Structure forms.
  • the resulting hot band coils which are a product of the hot rolling process is typically in the range of 1 to 20 mm in thickness.
  • Cold rolling is a widely used method for sheet production that is utilized to achieve targeted thickness for particular applications.
  • thinner gauges are usually targeted in the range of 0.4 to 2 mm.
  • cold rolling can be applied through multiple passes with or without intermediate annealing between passes. Typical reduction per pass is 5 to 70% depending on the material properties and equipment capability. The number of passes before the intermediate annealing also depends on materials properties and level of strain hardening during cold deformation.
  • the cold rolling will trigger Dynamic Nanophase Strengthening (Mechanism #2) leading to extensive strain hardening of the resultant sheet and to the formation of the High Strength Nanomodal Structure (Structure #3).
  • the properties of the cold rolled sheet from alloys herein will depend on the alloy chemistry and can be controlled by the cold rolling reduction to yield a fully cold rolled (i.e. hard) product or can be done to yield a range of properties (i.e. 1 ⁇ 4, 1 ⁇ 2, 3 ⁇ 4 hard etc.).
  • annealing is needed to recover the ductility of the material to allow for additional cold rolling gauge reduction.
  • Intermediate coils can be annealed by utilizing conventional methods such as batch annealing or continuous annealing lines.
  • the cold deformed High Strength Nanomodal Structure (Structure #3) for the steel alloys herein will undergo Recrystallization (Mechanism #3) during annealing leading to the formation of the Recrystallized Modal Structure (Structure #4).
  • the recrystallized coils can be a final product with advanced property combination depending on the alloy chemistry and targeted markets. In a case when even thinner gauges of the sheet are required, recrystallized coils can be subjected to further cold rolling to achieve targeted thickness that can be realized by one or multiple cycles of cold rolling / annealing.
  • the alloys herein are iron based metal alloys, having greater than or equal to 50 at.% Fe. More preferably, the alloys herein can be described as comprising, consisting essentially of, or consisting of the following elements at the indicated atomic percent: Fe (61.30 to 83.14 at. %); Si (0 to 7.02 at.%); Mn (0 to 15.86 at.%); B (0 to 6.09 at.%); Cr (0 to 18.90 at.%); Ni (0 to 8.68 at.%); Cu (0 to 2.00 at.%); C (0 to 3.72 at.%).
  • the alloys herein are such that they comprise Fe and at least four or more, or five or more, or six or more elements selected from Si, Mn, B, Cr, Ni, Cu or C. Most preferably, the alloys herein are such that they comprise, consist essentially of, or consist of Fe at a level of 50 at.% or greater along with Si, Mn, B, Cr, Ni, Cu and C.
  • Alloys were weighed out into charges ranging from 3,000 to 3,400 grams using commercially available ferroadditive powders with known chemistry and impurity content according to the atomic ratios in Table 2. Charges were loaded into a zirconia coated silica crucibles which was placed into an Indutherm VTC800V vacuum tilt casting machine. The machine then evacuated the casting and melting chambers and backfilled with argon to atmospheric pressure several times prior to casting to prevent oxidation of the melt. The melt was heated with a 14 kHz RF induction coil until fully molten, approximately 5.25 to 6.5 minutes depending on the alloy composition and charge mass.
  • Pre-heated slabs were pushed out of the tunnel furnace into a Fenn Model 061 2 high rolling mill.
  • the 50 mm slabs were preferably hot rolled for 5 to 8 passes though the mill before being allowed to air cool. After the initial passes each slab had been reduced between 80 to 85% to a final thickness of between 7.5 and 10 mm. After cooling each resultant sheet was sectioned and the bottom 190 mm was hot rolled for an additional 3 to 4 passes through the mill, further reducing the plate between 72 to 84% to a final thickness of between 1.6 and 2.1 mm.
  • Example pictures of laboratory cast slabs from two different alloys after hot rolling are shown in FIG. 4.
  • Hot rolling resultant sheets were media blasted with aluminum oxide to remove the mill scale and were then cold rolled on a Fenn Model 061 2 high rolling mill.
  • Cold rolling takes multiple passes to reduce the thickness of the sheet to a targeted thickness of typically 1.2 mm.
  • Hot rolled sheets were fed into the mill at steadily decreasing roll gaps until the minimum gap is reached. If the material has not yet hit the gauge target, additional passes at the minimum gap were used until 1.2 mm thickness was achieved. A large number of passes were applied due to limitations of laboratory mill capability.
  • Example pictures of cold rolled sheets from two different alloys are shown in FIG. 5.
  • Annealing la, lb, 2b were conducted in a Lucifer 7HT-K12 box furnace.
  • Annealing 2a and 3 was conducted in a Cameo Model G-ATM-12FL furnace. Specimens which were air normalized were removed from the furnace at the end of the cycle and allowed to cool to room temperature in air. For the furnace cooled specimens, at the end of the annealing the furnace was shut off to allow the sample to cool with the furnace. Note that the heat treatments were selected for demonstration but were not intended to be limiting in scope. High temperature treatments up to just below the melting points for each alloy are possible.
  • melting occurs in one or multiple stages with initial melting from ⁇ 1111°C depending on alloy chemistry and final melting temperature up to ⁇ 1476°C (Table 4). Variations in melting behavior reflect complex phase formation at solidification of the alloys depending on their chemistry.
  • the density of the alloys was measured on 9 mm thick sections of hot rolled material using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water.
  • the density of each alloy is tabulated in Table 5 and was found to be in the range from 7.57 to 7.89 g/cm 3 .
  • the accuracy of this technique is ⁇ 0.01 g/cm 3 .
  • Tensile properties were measured on an Instron 3369 mechanical testing frame using Instron' s Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron' s Advanced Video Extensometer. Tensile properties of the alloys listed in Table 2 after annealing with parameters listed in Table 3 are shown below in Table 6 to Table 10. The ultimate tensile strength values may vary from 799 to 1683 MPa with tensile elongation from 6.6 to 86.7%. The yield strength is in a range from 197 to 978 MPa. The mechanical characteristic values in the steel alloys herein will depend on alloy chemistry and processing conditions. The variation in heat treatment additionally illustrates the property variations possible through processing a particular alloy chemistry.
  • a laboratory slab with thickness of 50 mm was cast from Alloy 1 that was then laboratory processed by hot rolling, cold rolling and annealing at 850°C for 5 min as described in Main Body section of current application. Microstructure of the alloy was examined at each step of processing by SEM, TEM and x-ray analysis.
  • the cross section of the slab samples was ground on SiC abrasive papers with reduced grit size, and then polished progressively with diamond media paste down to 1 ⁇ . The final polishing was done with 0.02 ⁇ grit S1O 2 solution.
  • Microstructures were examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. To prepare TEM specimens, the samples were first cut by EDM, and then thinned by grinding with pads of reduced grit size every time. Further thinning to make foils of 60 to 70 ⁇ thickness was done by polishing with 9 ⁇ , 3 ⁇ and 1 ⁇ diamond suspension solution respectively.
  • Discs of 3 mm in diameter were punched from the foils and the final polishing was completed with electropolishing using a twin-jet polisher.
  • the chemical solution used was a 30% nitric acid mixed in methanol base.
  • the TEM specimens may be ion-milled using a Gatan Precision Ion Polishing System (PIPS).
  • PIPS Gatan Precision Ion Polishing System
  • the ion-milling usually is done at 4.5 keV, and the inclination angle is 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.
  • X-ray diffraction was done using a PANalytical X'Pert MPD diffractometer with a Cu Koc x-ray tube and operated at 45 kV with a filament current of 40 niA. Scans were run with a step size of 0.01° and from 25° to 95° two-theta with silicon incorporated to adjust for instrument zero angle shift. The resulting scans were then subsequently analyzed using Rietveld analysis using Siroquant software. Modal Structure was formed in the Alloy 1 slab with 50 mm thickness after solidification. The Modal Structure (Structure #1) is represented by a dendritic structure that is composed of several phases. In FIG.
  • the backscattered SEM image shows the dendritic arms that are shown in dark contrast while the matrix phase is in bright contrast. Note that small casting pores are found as exhibited (black holes) in the SEM micrograph.
  • TEM studies show that the matrix phase is primarily austenite (gamma-Fe) with stacking faults (FIG. 6b). The presence of stacking faults indicates a face-centered-cubic structure (austenite).
  • TEM also suggests that other phases could be formed in the Modal Structure.
  • a dark phase is found that identified as a ferrite phase with body-centered cubic structure (alpha-Fe) according to selected electron diffraction pattern.
  • FIG. 8a shows the backscattered SEM micrograph of Alloy 1 after being hot rolled from 50 mm to -1.7 mm at 1250°C. It can be seen that blocks of tens of microns in size are resulted from the dynamic recrystallization during the hot rolling, and the interior of the grains is relatively smooth indicating less amount of defects.
  • TEM further reveals that sub-grains of less than several hundred nanometers in size are formed, as shown in FIG. 8b.
  • X-ray diffraction analysis shows that the Nanomodal Structure of the Alloy 1 after hot rolling contains mainly austenite, with other phases such as ferrite and the iron manganese compound as shown in FIG. 9 and Table 12 .
  • FIG. 10a shows the backscattered SEM micrograph of cold rolled Alloy 1.
  • FIG. 10a shows the TEM micrograph of the microstructure in cold rolled Alloy 1. It can be seen that in addition to dislocations generated by the deformation, refined grains due to phase transformation can also be found.
  • the banded structure is related to the deformation twins caused by the cold rolling, corresponding to these in FIG 10a.
  • X-ray diffraction shows that the High Strength Nanomodal Structure of the Alloy 1 after cold rolling contains a significant amount of ferrite phase in addition to the retained austenite and the iron manganese compound as shown in FIG. 11 and Table .
  • Recrystallization occurs upon heat treatment of the cold deformed Alloy 1 with High Strength Nanomodal Structure (Structure #3,FIG. 1A and IB) that transforms into Recrystallized Modal Structure (Structure #4,FIG. IB).
  • the TEM images of the Alloy 1 after annealing are shown in ,FIG. 12.
  • equiaxed grains with sharp and straight boundaries are present in the structure and the grains are free of dislocations, which is characteristic feature of recrystallization.
  • the size of recrystallized grains can range from 0.5 to 50 ⁇ .
  • austenite is the dominant phase after recrystallization.
  • Annealing twins are occasionally found in the grains, but stacking faults are most often seen.
  • the formation of stacking faults shown in the TEM image is typical for face-centered-cubic crystal structure of austenite.
  • Backscattered SEM micrographs in FIG. 13 show the equiaxed recrystallized grains with the size of less than 10 ⁇ , consistent with TEM.
  • the different contrast of grains (dark or bright) seen on SEM images suggests that the crystal orientation of the grains is random, since the contrast in this case is mainly originated from the grain orientation. As a result, any texture formed by the previous cold deformation is eliminated.
  • X-ray diffraction shows that the Recrystallized Modal Structure of the Alloy 1 after annealing contains primarily austenite phase, with a small amount of ferrite and the iron manganese compound as shown in FIG. 14 and Table 14.
  • FIG. 15 shows the bright-field TEM micrographs of the microstructure in the deformed Alloy 1. Compared to the matrix grains that were initially almost dislocation-free in the Recrystallized Modal Structure after annealing, the application of stress generates a high density of dislocations within the matrix grains.
  • FIG. 15a shows local "pocket" of the transformed refined microstructure and selected area electron diffraction pattern corresponds to ferrite. Structural transformation into Refined High Strength Nanomodal Structure (Structure #5, FIG.
  • FIG. 16 shows the backscattered SEM images of the Refined High Strength Nanomodal Structure. Compared to the Recrystallized Modal Structure, the boundaries of matrix grains become less apparent, and the matrix is obviously deformed. Although the details of deformed grains cannot be revealed by SEM, the change caused by the deformation is enormous compared to the Recrystallized Modal Structure that was demonstrated in TEM images. X-ray diffraction shows that the Refined High Strength Nanomodal Structure of the Alloy 1 after tensile deformation contains a significant amount of ferrite and austenite phases.
  • the cross section of the slab samples was ground on SiC abrasive papers with reduced grit size, and then polished progressively with diamond media paste down to 1 ⁇ . The final polishing was done with 0.02 ⁇ grit S1O 2 solution.
  • Microstructures were examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. To prepare TEM specimens, the samples were first cut with EDM, and then thinned by grinding with pads of reduced grit size every time. Further thinning to make foils to -60 ⁇ thickness was done by polishing with 9 ⁇ , 3 ⁇ and 1 ⁇ 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 methanol base.
  • the TEM specimens may be ion-milled using a Gatan Precision Ion Polishing System (PIPS).
  • PIPS Gatan Precision Ion Polishing System
  • the ion-milling usually is done at 4.5 keV, and the inclination angle is 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.
  • X-ray diffraction was done using a Panalytical X'Pert MPD diffractometer with a Cu Koc x-ray tube and operated at 45 kV with a filament current of 40 mA. Scans were run with a step size of 0.01° and from 25° to 95° two- theta with silicon incorporated to adjust for instrument zero angle shift. The resulting scans were then subsequently analyzed using Rietveld analysis using Siroquant software.
  • FIG. 1A Modal Structure
  • Alloy 2 slab cast at 50 mm thick which is characterized by dendritic structure. Due to the presence of a boride phase (M 2 B), the dendritic structure is more evident than in Alloy 1 where borides are absent.
  • FIG. 18a shows the backscattered SEM of Modal Structure that exhibits a dendritic matrix (in bright contrast) with borides at the boundary (in dark contrast).
  • TEM studies show that the matrix phase is composed of austenite (gamma-Fe) with stacking faults (FIG. 18b). Similar to Alloy 1, the presence of stacking faults indicates the matrix phase is austenite.
  • FIG. 18b Also shown in TEM is the boride phase that appears dark in. FIG. 18b at the boundary of austenite matrix phase.
  • X-ray diffraction analysis data in. FIG. 19 and Table 16 shows that the Modal Structure contains austenite, M 2 B, ferrite, and iron manganese compound. Similar to Alloy 1, austenite is the dominant phase in the Alloy 2 Modal Structure, but other phases may be present depending on alloy chemistry.
  • FIG. 20a shows the backscattered SEM micrograph of hot rolled Alloy 2. Similar to Alloy 1, the dendritic Modal Structure is homogenized while the boride phase is randomly distributed in the matrix. TEM shows that the matrix phase is partially recrystallized as a result of dynamic recrystallization during hot rolling, as shown in FIG. 20b.
  • the matrix grains are on the order of 500 nm, which is finer than in Alloy 1 due to the pinning effect of borides.
  • X-ray diffraction analysis shows that the Nanomodal Structure of Alloy 2 after hot rolling contains mainly austenite phase and M2B, with other phases such as ferrite and iron manganese compound as shown in FIG. 21 and Table 17.
  • Table 17 X-ray Diffraction Data for Alloy 2 After Hot Rolling
  • FIG. 22a shows the backscattered SEM micrograph of the microstructure in the cold rolled Alloy 2. Deformation is concentrated in the matrix phase around the boride phase.
  • Recrystallization occurs upon annealing of the cold deformed Alloy 2 with High Strength Nanomodal Structure (Structure #3, FIG. 1A and IB) that transforms into Recrystallized Modal Structure (Structure #4, FIG. IB).
  • the recrystallized microstructure of the Alloy 2 after annealing is shown by TEM images in FIG. 24. As it can be seen, equiaxed grains with sharp and straight boundaries are present in the structure and the grains are free of dislocations, which is a characteristic feature of recrystallization.
  • the size of recrystallized grains is generally less than 5 ⁇ due to the pinning effect of boride phase, but larger grains are possible at higher annealing temperatures.
  • austenite is the dominant phase after recrystallization and stacking faults are present in the austenite, as shown in FIG. 24b.
  • the formation of stacking faults also indicates formation of face-centered-cubic austenite phase.
  • Backscattered SEM micrographs in FIG. 25 show the equiaxed recrystallized grains with the size of less than 5 ⁇ , with boride phase randomly distributed.
  • the different contrast of grains (dark or bright) seen on SEM images suggests that the crystal orientation of the grains is random, since the contrast in this case is mainly originated from the grain orientation. As a result, any texture formed by the previous cold deformation is eliminated.
  • X-ray diffraction shows that the Recrystallized Modal Structure of the Alloy 2 after annealing contains primarily austenite phase, with M2B, a small amount of ferrite, and a hexagonal phase with space group #186 (P6 3mc ) as shown in FIG. 26 and Table 19.
  • FIG. 27 shows the micrographs of microstructure in the deformed Alloy 2.
  • the initially dislocation-free matrix grains in the Recrystallized Modal Structure after annealing are filled with a high density of dislocations upon the application of stress, and the accumulation of dislocations in some grains activates the phase transformation from austenite to ferrite, leading to substantial refinement.
  • refined grains of 100 to 300 nm in size are shown in a local "pocket” where transformation occurred from austenite to ferrite.
  • Structural transformation into Refined High Strength Nanomodal Structure (Structure #5, FIG IB) in the "pockets" of matrix grains is a characteristic feature of the steel alloys herein.
  • FIG. 27b shows the backscattered SEM images of the Refined High Strength Nanomodal Structure.
  • Alloys were weighed out into charges ranging from 3,000 to 3,400 grams using commercially available ferroadditive powders with known chemistry and impurity content according to the atomic ratios in Table 2.
  • Charges were loaded into zirconia coated silica crucibles which were placed into an Indutherm VTC800V vacuum tilt casting machine. The machine then evacuated the casting and melting chambers and backfilled with argon to atmospheric pressure several times prior to casting to prevent oxidation of the melt.
  • the melt was heated with a 14 kHz RF induction coil until fully molten, approximately 5.25 to 6.5 minutes depending on the alloy composition and charge mass. After the last solids were observed to melt it was allowed to heat for an additional 30 to 45 seconds to provide superheat and ensure melt homogeneity.
  • the casting machine then evacuated the melting and casting chambers and tilted the crucible and poured the melt into a 50 mm thick, 75 to 80 mm wide, and 125 mm deep channel in a water cooled copper die.
  • the melt was allowed to cool under vacuum for 200 seconds before the chamber was filled with argon to atmospheric pressure.
  • Tensile specimens were cut from as-cast slabs by wire EDM and tested in tension. Results of tensile testing are shown in Table 21. As it can be seen, ultimate tensile strength of the alloys herein in as-cast condition varies from 411 to 907 MPa. The tensile elongation varies from 3.7 to 24.4%. Yield strength is measured in a range from 144 to 514 MPa.
  • resultant sheets were media blasted with aluminum oxide to remove the mill scale and were then cold rolled on a Fenn Model 061 2 high rolling mill.
  • Cold rolling takes multiple passes to reduce the thickness of the sheet to targeted thickness, generally 1.2 mm.
  • Hot rolled sheets were fed into the mill at steadily decreasing roll gaps until the minimum gap is reached. If the material has not yet hit the gauge target, additional passes at the minimum gap were used until the targeted thickness was reached.
  • Cold rolling conditions with the number of passes for each alloy herein are listed in Table 23.
  • Tensile specimens were cut from cold rolled sheets by wire EDM and tested in tension. Results of tensile testing are shown in Table 23. Cold rolling leads to significant strengthening with ultimate tensile strength in the range from 1356 to 1831 MPa.
  • the tensile elongation of the alloys herein in cold rolled state varies from 1.6 to 32.1%. Yield strength is measured in a range from 793 to 1645 MPa. It is anticipated that higher ultimate tensile strength and yield strength can be achieved in alloys herein by larger cold rolling reduction (>40%) that in our case is limited by laboratory mill capability. With more rolling force, it is anticipated that ultimate tensile strength could be increased to at least 2000 MPa and yield strength to at least 1800 MPa.
  • Hot rolled sheet from each alloy was subjected to three cycles of cold rolling and annealing. Sheet thicknesses before and after hot rolling and cold rolling reduction at each cycle are listed in Table 25. Annealing at 850°C for 10 min was applied after each cold rolling. Tensile specimens were cut from the sheet in the initial hot rolled state and at each step of the cycling. Tensile properties were measured on an Instron 3369 mechanical testing frame using Instron' s Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron' s Advanced Video Extensometer.
  • Alloy 1 has ultimate tensile strength from 1216 to 1238 MPa in hot rolled state with ductility from 50.0 to 52.7% and yield strength from 264 to 285 MPa.
  • the ultimate tensile strength was measured in the range from 1482 to 1517 MPa at each cycle.
  • Ductility was found consistently in the range from 28.5 to 32.8% with significantly higher yield strength of 718 to 830 MPa as compared to that in hot rolled condition.
  • Annealing at each cycle resulted in restoration of the ductility to the range from 47.7 to 59.7% with ultimate tensile strength from 1216 to 1270 MPa.
  • Yield strength after cold rolling and annealing is lower than that after cold rolling and was measured in the range from 431 to 515 MPa that is however higher than that in initial hot rolled condition. Similar results with property reversibility between cold rolled and annealed material through cycling were observed for Alloy 2 (FIG. 32b). In initial hot rolled state, Alloy 2 has ultimate tensile strength from 1219 to 1277 MPa with ductility from 41.9 to 48.2% and yield strength from 454 to 480 MPa. Cold rolling at each cycle results in the material strengthening to the ultimate tensile strength from 1553 to 1598 MPa with ductility reduction to the range from 20.3 to 24.1%. Yield strength was measured from 912 to 1126 MPa.
  • Alloy 2 After annealing at each cycle, Alloy 2 has ultimate tensile strength from 1231 to 1281 MPa with ductility from 46.9 to 53.5%. Yield strength in Alloy 2 after cold rolling and annealing at each cycle is similar to that in hot rolled condition and varies from 454 to 521 MPa.
  • Bend tests were performed using an Instron 5984 tensile test platform with an Instron W- 6810 guided bend test fixture according to specifications outlined in the ISO 7438 International Standard Metallic materials— Bend test (International Organization for Standardization, 2005). Test specimens were cut by wire EDM to a dimension of 20 mm x 55 mm x sheet thickness. No special edge preparation was done to the samples. Bend tests were performed using an Instron 5984 tensile test platform with an Instron W-6810 guided bend test fixture. Bend tests were performed according to specifications outlined in the ISO 7438 International Standard Metallic materials— Bend test (International Organization for Standardization, 2005).
  • the test was performed by placing the test specimen on the fixture supports and pushing with a former as shown in FIG. 33.
  • the specimens Prior to bending, the specimens were lubricated on both sides with 3 in 1 oil to reduce friction with the test fixture. This test was performed with a 1 mm diameter former. The former was pushed downward in the middle of the supports to different angles up to 180° or until a crack appeared. The bending force was applied slowly to permit free plastic flow of the material. The displacement rate was calculated based on the span gap of each test in order to have a constant angular rate and applied accordingly.
  • Results of the bending response of the alloys herein are listed in Table 28 including initial sheet thickness, former radius to sheet thickness ratio (r/t) and maximum bend angle before cracking. All alloys listed in the Table 28 did not show cracks at 90° bend angle. The majority of the alloys herein have capability to be bent at 180° angle without cracking. Example of the samples from Alloy 1 after bend testing to 180° is shown in FIG. 34.
  • shearing may occur herein by a number of processing options, such as piercing, perforating, cutting or cropping (cutting off of an end of a given metal part).
  • Tensile specimens in the ASTM E8 geometry were prepared using both wire EDM cutting and punching. Tensile properties were measured on an Instron 5984 mechanical testing frame using Instron' s Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron' s Advanced Video Extensometer. Tensile data is shown in Table 29 and illustrated in FIG. 35a for selected alloys. Decrease in properties is observed for all alloys tested but the level of this decrease varies significantly depending on alloy chemistry. Table 30 summarizes a comparison of ductility in punched samples as compared to that in the wire EDM cut samples. In FIG.
  • EDM cutting is considered to be representative of the optimal mechanical properties of the identified alloys, without a sheared edge, and which were processed to the point of assuming Structure #4 (Recrystallized Modal Structure). Accordingly, samples having a sheared edge due to punching indicate a significant drop in ductility as reflected by tensile elongation measurements of the punched samples having the ASTM E8 geometry. For Alloy 1, tensile elongation is initially 47.2% and then drops to 8.1%, a drop itself of 82.8%%. The drop in ductility from the punched to the EDM cut (E2/E1) varies from 0.57 to 0.05.
  • the edge status after punching and EDM cutting was analyzed by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.
  • the typical appearance of the specimen edge after EDM cutting is shown for Alloy 1 in FIG. 36a.
  • the EDM cutting method minimizes the damage of a cut edge allowing the tensile properties of the material to be measured without any deleterious edge effects.
  • wire-EDM cutting material is removed from the edge by a series of rapidly recurring current discharges / sparks and by this route an edge is formed without substantial deformation or edge damage.
  • the appearance of the sheared edge after punching is shown in FIG. 36b.
  • annealing may be achieved by various methods, including but not limited to furnace heat treatment, induction heat treatment and/or laser heat treatment.
  • Tensile specimens in the ASTM E8 geometry were prepared using both wire EDM cutting and punching. Part of punched tensile specimens was then put through a recovery anneal of 850°C for 10 minutes, followed by an air cool, to confirm the ability to recover properties lost by punching and shearing damage. Tensile properties were measured on an Instron 5984 mechanical testing frame using Instron' s Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron' s Advanced Video Extensometer. Tensile testing results are provided in Table 31 and illustrated in FIG. 38 for selected alloys showing a substantial mechanical property recovery in punched samples after annealing.
  • a tensile elongation average value is about 47.2%.
  • the tensile testing of the sample with such edge indicated a significant drop in such elongation values, i.e. an average value of only about 8.1% due to Mechanism #4 and formation of Refined High Strength Nanomodal Structure (Structure #5, FIG. IB), which while present largely at the edge section where shearing occurred, is nonetheless reflected in the bulk property measurements in tensile testing.
  • Mechanism #3 which is representative of Mechanism #3 in FIG.
  • tensile strength is initially at an average of about 47.1%, dropping to an average value of 8.1%, a decrease of up to about 80 to 85%, and upon annealing and undergoing what is shown in FIG. IB as Mechanism #3, tensile elongation recovers to an average value of 46.1%, a recovery of greater than or equal to 90% of the value of the elongation value of 47.1%.
  • FIG. 40 Tensile testing results are shown in FIG. 40 demonstrating a transition in deformation behavior depending on annealing temperature.
  • the Dynamic Nanophase Strengthening (Mechanism #2, FIG. 1A) or the Nanophase Refinement & Strengthening (Mechanism #4, FIG. IB) occurs which involves, once the yield strength is exceeded with increasing strain, the continuous transformation of austenite to ferrite plus one or more types of nanoscale hexagonal phases. Concurrent with this transformation, deformation by dislocation mechanisms also occurs in the matrix grains prior to and after transformation. The result is the change in the microstructure from the Nanomodal Structure (Structure #2, FIG. 1A) to the High Strength Nanomodal Structure (Structure #3, FIG.
  • the yield strength can be varied widely from 1372 MPa at the 500°C anneal down to 458 MPa at the 850°C anneal.
  • FIG. 41 shows the microstructure of as-cold rolled Alloy 1 sample.
  • Annealing at 650°C for 10 min shows larger recrystallized grains suggesting the progress of recrystallization. Although the fraction of deformed area is reduced, the deformed structure continues to be seen, as shown in FIG. 44. Annealing at 700°C 10 min shows larger and cleaner recrystallized grains, as displayed by FIG. 45. Selected electron diffraction shows that these recrystallized grains are of the austenite phase. The area of deformed structure is smaller compared to the samples annealed at lower temperature. Survey over the entire sample suggests that approx. 10% to 20% area is occupied by the deformed structure. The progress of recrystallization revealed by TEM in the samples annealed at lower temperature to higher temperature corresponds excellently to the change of tensile properties shown in FIG. 40.
  • TTT diagram in FIG. 46 One reason behind the difference in recovery and transition in deformation behavior is illustrated by the model TTT diagram in FIG. 46. As described previously, the very fine / nanoscale grains of ferrite formed during cold working recrystallize into austenite during annealing and some fraction of the nanoprecipitates re-dissolve. Concurrently, the effect of the strain hardening is eliminated with dislocation networks and tangles, twin boundaries, and small angle boundaries being annihilated by various known mechanisms. As shown by the heating curve A of the model temperature, time transformation (TTT) diagram in FIG. 46, at low temperatures (particularly below 650°C for Alloy 1), only recovery may occur without recrystallization (i.e. recovery being a reference to a reduction in dislocation density).
  • TTT time transformation
  • the effect of shearing and formation of a sheared edge, and its associated negative influence on mechanical properties can be at least partially recovered at temperatures of 450°C up to 650°C as shown in FIG. 46.
  • recrystallization can occur, which also contributes to restoring mechanical strength lost due to the formation of a sheared edge.
  • this Case Example demonstrates that upon deformation during cold rolling, concurrent processes occur involving dynamic strain hardening and phase transformation through unique Mechanisms #2 or #3 (FIG. 1A) along with dislocation based mechanisms.
  • the microstructure Upon heating, the microstructure can be reversed into a Recrystallized Modal Structure (Structure #4, FIG. IB).
  • Structure #4 FIG. IB
  • this reversing process may not occur when only dislocation recovery takes place.
  • various external heat treatments can be used to heal the edge damage from punching / stamping.
  • Tensile specimens in the ASTM E8 geometry were prepared by punching. A part of punched tensile specimens from selected alloys was then put through a recovery anneal for 10 minutes at different temperatures in a range from 450 to 850°C, followed by an air cool. Tensile properties were measured on an Instron 5984 mechanical testing frame using Instron's Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron's Advanced Video Extensometer.
  • FIG. 49 shows the backscattered SEM images of the microstructure at the edge in the as- punched condition. It can be seen that the microstructure is deformed and transformed in the shear affected zone (i.e., the triangle with white contrast close to the edge) in contrast to the recrystallized microstructure in the area away from the shear affected zone. Similar to tensile deformation, the deformation in the shear affected zone caused by punching creates Refined High Strength Nanomodal Structure (Structure #5, FIG. IB) through Nanophase Refinement & Strengthening mechanism. However, annealing recovers the tensile properties of punched ASTM E8 specimens, which are related to the microstructure change in the shear affected zone during annealing.
  • FIG. 49 shows the backscattered SEM images of the microstructure at the edge in the as- punched condition.
  • Tensile specimens in the ASTM E8 geometry were prepared by punching at three different speeds of 28 mm/s, 114 mm/s, and 228 mm/s. Wire EDM cut specimens from the same materials were used for the reference. A part of punched tensile specimens from selected alloys was then put through a recovery anneal for 10 minutes at 850°C, followed by an air cool. Tensile properties were measured on an Instron 5984 mechanical testing frame using Instron's Bluehill control software. All tests were conducted at room temperature, with the bottom grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was collected using Instron's Advanced Video Extensometer.
  • Specimens for testing with a size of 89 x 89 mm were wire EDM cut from the sheet.
  • the hole with 10 mm diameter was cut in the middle of specimens by utilizing two methods: punching and drilling with edge milling.
  • the hole punching was done on an Instron Model 5985 Universal Testing System using a fixed speed of 0.25 mm/s with 16% clearance.
  • Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press and consisted of slowly raising the conical punch that uniformly expanded the hole radially outward. A digital image camera system was focused on the conical punch and the edge of the hole was monitored for evidence of crack formation and propagation.
  • the initial diameter of the hole was measured twice with calipers, measurements were taken at 90° increments and averaged to get the initial hole diameter.
  • the conical punch was raised continuously until a crack was observed propagating through the specimen thickness. At that point the test was stopped and the hole expansion ratio was calculated as a percentage of the initial hole diameter measured before the start of the test. After expansion four diameter measurements were taken using calipers every 45° and averaged to account for any asymmetry of the hole due to cracking.
  • Results of HER testing are shown in FIG. 53 demonstrating a significantly lower value for the sample when the hole was prepared by punching as compared to milling: 5.1% HER vs 73.6% HER, respectively. Samples were cut from both tested samples as shown in FIG. 54 for SEM analysis and microhardness measurements.
  • Microhardness was measured for Alloy 1 at all relevant stages of the hole expansion process. Microhardness measurements were taken along cross sections of sheet samples in the annealed (before punching and HER testing), as-punched, and HER tested conditions. Microhardness was also measured in cold rolled sheet from Alloy 1 for reference. Measurement profiles started at an 80 micron distance from the edge of the sample, with an additional measurement taken every 120 microns until 10 such measurements were taken. After that point, further measurements were taken every 500 microns, until at least 5 mm of total sample length had been measured. A schematic illustration of microhardness measurement locations in HER tested samples is shown in FIG. 55. SEM images of the punched and HER tested samples after microhardness measurements are shown in FIG. 56.
  • the punching process creates a transformed zone of approximately 500 microns immediately adjacent to the punched edge, with the material closest to the punched edge either fully or near-fully transformed, as evidenced by the hardness approaching that observed in the fully-transformed, 40% cold rolled material immediately next to the punched edge.
  • Microhardness profiles for each sample is presented in FIG. 58. As it can be seen, microhardness gradually increases towards a hole edge in the case of milled while in the case of punched hole microhardness increase was observed in a very narrow area close to the hole edge. TEM samples were cut at the same distance in both cases as indicated in FIG. 58.
  • the HER test samples were first sectioned by wire EDM, and a piece with a portion of hole edge was thinned by grinding with pads of reduced grit size. Further thinning to -60 ⁇ thickness is done by polishing with 9 ⁇ , 3 ⁇ , and 1 ⁇ diamond suspension solution respectively. Discs of 3 mm in diameter were punched from the foils near the edge of the hole and the final polishing was completed by electropolishing using a twin-jet polisher. The chemical solution used was a 30% Nitric acid mixed in Methanol base. In case of insufficient thin area for TEM observation, the TEM specimens may be ion-milled using a Gatan Precision Ion Polishing System (PIPS).
  • PIPS Gatan Precision Ion Polishing System
  • the ion-milling usually is done at 4.5 keV, and the inclination angle is 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. Since the location for TEM study is at the center of the disc, the observed microstructure is approximately -1.5 mm from the edge of hole.
  • FIB Focused Ion Beam
  • Test specimens of 89 x 89 mm were wire EDM cut from the sheet from larger sections.
  • a 10 mm diameter hole was made in the center of specimens by punching on an Instron Model 5985 Universal Testing System using a fixed speed of 0.25 mm/s at 16% punch to die clearance.
  • Half of the prepared specimens with punched holes were individually wrapped in stainless steel foil and annealed at 850°C for 10 minutes before HER testing.
  • Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press and consisted of slowly raising the conical punch that uniformly expanded the hole radially outward.
  • a digital image camera system was focused on the conical punch and the edge of the hole was monitored for evidence of crack formation and propagation.
  • the initial diameter of the hole was measured twice with calipers, measurements were taken at 90° increments and averaged to get the initial hole diameter.
  • the conical punch was raised continuously until a crack was observed propagating through the specimen thickness. At that point the test was stopped and the hole expansion ratio was calculated as a percentage of the initial hole diameter measured before the start of the test. After expansion four diameter measurements were taken using calipers every 45° and averaged to account for any asymmetry of the hole due to cracking.
  • the results of the hole expansion ratio measurements on the specimens with and without annealing after hole punching are shown in Table 35.
  • the hole expansion ratio measured with punched holes with annealing is generally greater than in punched holes without annealing.
  • the increase in hole expansion ratio with annealing for the identified alloys herein therefore leads to an increase in the actual HER of about 25% to 90%.

Abstract

L'invention concerne des procédés d'amélioration des propriétés mécaniques d'un alliage métallique ayant subi une ou plusieurs pertes de propriétés mécaniques suite à un cisaillement, par exemple dans la formation d'une partie bord cisaillée ou d'un trou poinçonné. Les procédés selon l'invention permettent d'améliorer les propriétés mécaniques d'alliages métalliques ayant été formés avec un ou plusieurs bords cisaillés qui peuvent sinon s'avérer être un facteur limitant pour des applications industrielles.
PCT/US2018/018751 2017-02-21 2018-02-20 Formabilité de bord améliorée dans les alliages métalliques WO2018160387A1 (fr)

Priority Applications (6)

Application Number Priority Date Filing Date Title
CA3053383A CA3053383A1 (fr) 2017-02-21 2018-02-20 Formabilite de bord amelioree dans les alliages metalliques
EP18760610.8A EP3585532A4 (fr) 2017-02-21 2018-02-20 Formabilité de bord améliorée dans les alliages métalliques
JP2019544895A JP2020509233A (ja) 2017-02-21 2018-02-20 金属合金における端部形成能の改善
CN201880013756.6A CN110382130A (zh) 2017-02-21 2018-02-20 金属合金的改善的边缘可成形性
KR1020197027312A KR20190130131A (ko) 2017-02-21 2018-02-20 금속 합금의 개선된 에지 성형성
JP2023039718A JP2023075277A (ja) 2017-02-21 2023-03-14 金属合金における端部形成能の改善

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
US15/438,313 2017-02-21
US15/438,313 US10465260B2 (en) 2015-04-10 2017-02-21 Edge formability in metallic alloys

Publications (1)

Publication Number Publication Date
WO2018160387A1 true WO2018160387A1 (fr) 2018-09-07

Family

ID=63370238

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/US2018/018751 WO2018160387A1 (fr) 2017-02-21 2018-02-20 Formabilité de bord améliorée dans les alliages métalliques

Country Status (6)

Country Link
EP (1) EP3585532A4 (fr)
JP (2) JP2020509233A (fr)
KR (1) KR20190130131A (fr)
CN (1) CN110382130A (fr)
CA (1) CA3053383A1 (fr)
WO (1) WO2018160387A1 (fr)

Families Citing this family (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN111312439A (zh) * 2020-04-01 2020-06-19 吉林省中赢高科技有限公司 一种电能传输铝件及其加工工艺

Citations (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2013167572A1 (fr) * 2012-05-08 2013-11-14 Tata Steel Ijmuiden Bv Pièce de châssis automobile en tôle d'acier laminé à chaud très résistante
US20160303635A1 (en) * 2015-04-10 2016-10-20 The Nanosteel Company, Inc. Edge Formability In Metallic Alloys

Family Cites Families (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5470529A (en) * 1994-03-08 1995-11-28 Sumitomo Metal Industries, Ltd. High tensile strength steel sheet having improved formability
CA2387322C (fr) * 2001-06-06 2008-09-30 Kawasaki Steel Corporation Tole d'acier a ductilite elevee possedant des proprietes superieures de formabilite sous pressage et de vieillissement par ecrouissage, et methode de fabrication dudit produit
JP4473588B2 (ja) * 2004-01-14 2010-06-02 新日本製鐵株式会社 めっき密着性および穴拡げ性に優れた溶融亜鉛めっき高強度鋼板の製造方法
JP4510488B2 (ja) * 2004-03-11 2010-07-21 新日本製鐵株式会社 成形性および穴拡げ性に優れた溶融亜鉛めっき複合高強度鋼板およびその製造方法
JP4445365B2 (ja) * 2004-10-06 2010-04-07 新日本製鐵株式会社 伸びと穴拡げ性に優れた高強度薄鋼板の製造方法
JP5136609B2 (ja) * 2010-07-29 2013-02-06 Ｊｆｅスチール株式会社 成形性および耐衝撃性に優れた高強度溶融亜鉛めっき鋼板およびその製造方法
CN108456832B (zh) * 2012-02-27 2021-02-02 株式会社神户制钢所 弯曲加工性优良的超高强度冷轧钢板及其制造方法
JP6626436B2 (ja) * 2013-10-02 2019-12-25 ザ・ナノスティール・カンパニー・インコーポレーテッド 先進的高強度金属合金の生産のための再結晶、微細化、及び強化機構

Patent Citations (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2013167572A1 (fr) * 2012-05-08 2013-11-14 Tata Steel Ijmuiden Bv Pièce de châssis automobile en tôle d'acier laminé à chaud très résistante
US20160303635A1 (en) * 2015-04-10 2016-10-20 The Nanosteel Company, Inc. Edge Formability In Metallic Alloys

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
TAYLOR, MD: "Effect of microstructure on the fracture response of advanced high strength steels", THESIS . COLORADO SCHOOL OF MINES, 2016, XP055537697, Retrieved from the Internet <URL:https://mountainscholar.org/bitstream/handle/11124/170014/Taylor_mines_0052E_10913.pdf?sequence=1> [retrieved on 20180321] *

Also Published As

Publication number Publication date
JP2020509233A (ja) 2020-03-26
CA3053383A1 (fr) 2018-09-07
EP3585532A4 (fr) 2020-08-12
KR20190130131A (ko) 2019-11-21
CN110382130A (zh) 2019-10-25
JP2023075277A (ja) 2023-05-30
EP3585532A1 (fr) 2020-01-01

Similar Documents

Publication Publication Date Title
CA2982346C (fr) Aptitude au formage de bord amelioree dans des alliages metalliques
KR102256921B1 (ko) 첨단 고강도 금속 합금의 제조를 위한 재결정화, 미세화, 및 강화 메커니즘
US8257512B1 (en) Classes of modal structured steel with static refinement and dynamic strengthening and method of making thereof
KR102012956B1 (ko) 고 강도 및 고 연성을 갖는 비-스테인리스 강의 신규 부류
Kumar et al. Enhancement of mechanical properties of low stacking fault energy brass processed by cryorolling followed by short-annealing
US20210238703A1 (en) Edge formability in metallic alloys
EP3481972A1 (fr) Acier à haute limite d&#39;élasticité
US11254996B2 (en) Delayed cracking prevention during drawing of high strength steel
JP2023075277A (ja) 金属合金における端部形成能の改善
EP3645763A1 (fr) Rétention de propriétés mécaniques dans des alliages d&#39;acier après traitement et en présence de sites de concentration de contraintes

Legal Events

Date Code Title Description
121 Ep: the epo has been informed by wipo that ep was designated in this application

Ref document number: 18760610

Country of ref document: EP

Kind code of ref document: A1

ENP Entry into the national phase

Ref document number: 3053383

Country of ref document: CA

ENP Entry into the national phase

Ref document number: 2019544895

Country of ref document: JP

Kind code of ref document: A

NENP Non-entry into the national phase

Ref country code: DE

ENP Entry into the national phase

Ref document number: 20197027312

Country of ref document: KR

Kind code of ref document: A

ENP Entry into the national phase

Ref document number: 2018760610

Country of ref document: EP

Effective date: 20190923