KR20190130131A - Improved edge formability of metal alloys - Google Patents

Abstract

The present disclosure is directed to a method for improving the mechanical properties of a metal alloy that suffers a loss of one or more mechanical properties as a result of forming the edge, such as the formation of an inner hole or an outer edge. A method is disclosed that provides the ability to improve the mechanical properties of a metal alloy formed with one or more edges disposed within the metal alloy by various methods that can serve as limiting factors for industrial applications.

Description

Improved edge formability of metal alloys

Cross Reference of Related Application

This application is directed to US Provisional Patent Application Serial No. 62 / 146,048, filed April 10, 2015, which is incorporated herein in its entirety by reference, and US Provisional Patent Application Serial No. 62 / 257,070, filed November 18, 2015. A priority claim is U.S. Patent Application 15 / 438,313, filed February 21, 2017, which is a partial continuing application of U.S. Patent Application 15 / 094,554, filed April 8, 2016, claiming benefit.

Field of invention

The present disclosure is directed to a method of improving the mechanical properties of a metal alloy that has lost one or more mechanical properties as a result of shearing, such as sheared edge portions or punched holes. More specifically, a method is disclosed that provides the ability to improve the mechanical properties of a metal alloy formed with one or more sheared edges that can serve as a limiting factor in industrial applications.

From ancient tools to modern skyscrapers and automobiles, the river has been driving human innovation for hundreds of years. Iron and its associated alloys, rich in the earth's crust, have provided mankind with solutions to many difficult developmental barriers. Since its inception, river development has advanced considerably within the last two centuries, with new types of steel being introduced every few years. These steel alloys can be classified into three classes based on measured properties, in particular maximum tensile strain and pre-destructive tensile stress. These three classes are Low Strength Steel (LSS), High Strength Steel (HSS), and Advanced High Strength Steel (AHSS). Low strength steels (LSS) are generally classified as exhibiting a maximum tensile strength of less than 270 MPa and include types such as interstitial free steels and mild steels. High strength steels (HSS) are classified as exhibiting maximum tensile strengths of 270 to 700 MPa and include types such as high strength low alloy steels, high strength interstitial free steels and heat curable steels. Ultra high strength steels (AHSS) are classified by maximum tensile strengths in excess of 700 MPa, such as martensitic steel (MS), double phase (DP) steel, transformation organic plastic (TRIP) steel, and composite phase (CP) steel. Include type. As the level of strength increases, the tendency of the maximum tensile elongation (melt) of the steel becomes negative, and the elongation decreases at high maximum tensile strength. For example, the tensile elongations of LSS, HSS and AHSS range from 25% to 55%, 10% to 45%, and 4% to 30%, respectively.

Steel production continues to increase, with current US production of about 100 million tons per year, estimated at $ 75 billion. In-vehicle utilization is also high, ultra-high strength steel (AHSS) is currently 17%, and 300% growth is expected in the next few years [American Iron and Steel Institute. (2013). Profile 2013. Washington, D.C.]. Current market trends and government regulations are driving the efficiency of vehicles, and AHSS is increasingly pursuing the ability to provide high strength-to-mass ratios. The high strength of AHSS allows designers to reduce the thickness of finished parts while maintaining equivalent or improved mechanical properties. When reducing the thickness of the part, vehicle fuel efficiency is improved because less mass is needed to obtain equivalent or better mechanical properties of the vehicle. This allows the designer to improve the fuel efficiency of the vehicle without compromising safety.

One of the important properties of next generation steel is formability. Formability is the ability to make a specific shape without cracking, rupturing or otherwise breaking. High formability steels benefit component designers by enabling the creation of more complex part shapes and by reducing weight. Formability can be further classified into two distinct forms of edge formability and bulk formability. Edge formability is the ability of an edge to be formed into a particular shape. Edges of the material are produced through various methods of industrial processes, including but not limited to punching, shearing, piercing, stamping, perforating, cutting, or cropping. In addition, the apparatus used to create such edges is as varied as the methods including, but not limited to, various types of mechanical presses, hydraulic presses, and / or electromagnetic presses. Depending on the application and the material being processed, the range of edge generation rates also varies widely at low speeds of 0.25 mm / sec and high speeds of 3700 mm / sec. Various edge forming methods, devices, and speeds result in a myriad of different edge conditions that are commercially available today.

Edges, which are free surfaces, are governed by defects such as cracks or structural changes in the sheet resulting from the creation of the sheet edges. These defects adversely affect edge formability during molding operations, thereby reducing the effective malleability at the edges. Bulk formability, on the other hand, is governed by the inherent malleability, structure, and associated stress state of the metal during the molding operation. Bulk formability is mainly affected by available deformation mechanisms such as dislocations, twins, and phase transformations. Bulk formability is maximized when these available deformation mechanisms are saturated in the material, and improved bulk formability is obtained by increasing the number and availability of these mechanisms.

Edge formability can be measured through hole expansion measurements in which holes are formed in the sheet, which holes are expanded by a conical punch. Previous studies have shown that conventional AHSS materials suffer from reduced edge formability compared to 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]. For example, dual phase (DP) steel with a maximum tensile strength of 780 MPa achieves less than 20% hole expansion, while interstitial free steel (IF) with a maximum tensile strength of about 400 MPa is about 100% Achieve hole expansion rate. This reduced edge formability makes it difficult to employ AHSS in automotive applications, despite having desirable bulk formability.

summary

A method of improving one or more mechanical properties in a metal alloy that has suffered a loss of mechanical properties as a result of the formation of one or more sheared edges includes the following steps:

a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy having Tm and matrix grains of μm to 10,000 μm;

b. The alloy is heated to a temperature of 700 ° C. to less than the Tm of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , and the resulting first having a maximum tensile strength of 921 MPa to 1413 MPa. Providing an alloy;

c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;

d. Heating the resulting second alloy to a temperature below T _m and forming the resulting third alloy having a matrix grain and elongation (E ₁ ) of 0.5 μm to 50 μm;

e. Shearing the alloy to form one or more sheared edges, the elongation of the alloy being reduced to an E ₂ value, wherein E ₂ = (0.57 to 0.05) (E ₁ );

f. Reheating the alloy with the one or more sheared edges. Here, the reduced elongation of the alloys observed in step (d) is restored to the level of elongation E ₃ = (0.48 to 1.21) (E _1).

The present disclosure also relates to a method for improving the hole expansion rate in a metal alloy in which the hole expansion rate is lost as a result of forming a hole with a sheared edge, comprising the following steps:

a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy having Tm and matrix grains of μm to 10,000 μm;

b. The alloy is heated to a temperature of 700 ° C. to below the T _m of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , a maximum tensile strength of 921 MPa to 1413 MPa and 12.0% to 77.7 Providing a resultant first alloy having an elongation of%;

c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;

d. Heat the resulting second alloy to a temperature below 650 ° C. to Tm, form the resulting third alloy with matrix grains from 0.5 μm to 50 μm, and form a hole in the resulting third alloy by shearing. The hole has a sheared edge and has a first hole expansion ratio HER ₁ ; And

e. Heating the alloy with the hole and the associated HER _1. Here, the alloy represents the second hole expansion rate (HER ₂ ), where HER ₂ > HER ₁ .

The present invention also relates to a method for improving the hole expansion rate in a metal alloy in which the hole expansion rate is lost as a result of forming a hole with sheared edges, comprising the following steps:

a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy having a Tm and matrix grains of 2 μm to 10,000 μm;

b. The alloy is heated to a temperature of 700 ° C. to below the T _m of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , a maximum tensile strength of 921 MPa to 1413 MPa and 12.0% to 77.7 Providing a resultant first alloy having an elongation of%;

c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;

d. Heating the resulting second alloy to a temperature below 650 ° C. to Tm and forming the resulting third alloy having matrix grains of 0.5 μm to 50 μm, wherein the alloy is 30 to − for a hole formed without shear; Characterized in that it has a first hole expansion ratio HER ₁ of 130%;

e. Forming a hole in the resulting second alloy, the hole formed by shearing, exhibiting a second hole expansion ratio (HER ₂ ), wherein HER ₂ = (0.01 to 0.30) (HER ₁ );

f. Heating the alloy. Wherein HER ₂ is restored to a value of HER ₃ = (0.60 to 1.0) HER ₁ .

The invention also relates to a method of punching one or more holes in a metal alloy, comprising the following steps:

a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy having a Tm and matrix grains of 2 μm to 10,000 μm;

b. The alloy is heated to a temperature of 700 ° C. to below the T _m of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , a maximum tensile strength of 921 MPa to 1413 MPa and 12.0% to 77.7 Providing a resultant first alloy having an elongation of%;

c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;

d. Heating the resulting second alloy to a temperature below 650 ° C. to Tm and forming the resulting third alloy having a matrix grain and elongation (E ₁ ) of 0.5 μm to 50 μm;

e. Punching holes in the alloy at a punch speed of at least 10 mm / sec. Wherein the punched hole exhibits a hole expansion rate of at least 10%.

The invention also relates to a method of expanding an edge in an alloy, comprising the following steps:

a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy with Tm;

b. The alloy is heated to a temperature of 700 ° C. to less than the Tm of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , a maximum tensile strength of 921 MPa to 1413 MPa and 12.0% to 77.7% Providing a resultant first alloy having an elongation of;

c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;

d. Heating the resulting second alloy to a temperature below Tm and forming the resulting third alloy having an elongation of 6.6% to 86.7%;

e. Forming an edge in the resulting alloy, and extending the edge at a speed of at least 5 mm / min.

The invention also relates to a method of expanding the edge of an alloy, comprising the following steps:

Supplying a metal alloy comprising at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, the alloy having a maximum tensile strength of 799 MPa to 1683 MPa and Having an elongation of 6.6 to 86.7%;

Forming an edge in the alloy; And

Expanding the edges in the alloy at a speed of at least 5 mm / min.

The following detailed description may be better understood with reference to the accompanying drawings, which are provided for illustrative purposes and should not be construed as limiting any aspect of the invention.

1A is a structural pathway for the formation and associated mechanisms of high strength nanomodal structures.
1B is a structural pathway for the formation and related mechanisms of recrystallized modal structures and refined high strength nanomodal structures.
2 is a structural pathway for the development of refined high strength nanomodal structures related to industrial process steps.
3 is an image of a 50 mm slab laboratory cast with a) alloy 9 and b) alloy 12. FIG.
4 is an image of a hot rolled sheet after laboratory casting with a) Alloy 9 and b) Alloy 12. FIG.
FIG. 5 is an image of a cold rolled sheet after laboratory casting and hot rolling with a) Alloy 9 and b) Alloy 12.
6 is a microstructure of solidified alloy 1 cast to a thickness of 50 mm, a) is a backscatter SEM micrograph showing the dendritic properties of the modal structure in the cast state, and b) a bright field showing the details of the matrix grains TEM micrograph, c) bright field TEM with selected electron diffraction showing ferrite phase in the modal structure.
7 is an X-ray diffraction pattern of the modal structure of the alloy 1 alloy after solidification, a) is experimental data, and b) Rietveld micronization analysis.
8 is a microstructure of Alloy 1 after hot rolling to 1.7 mm thickness, a) is backscatter SEM micrograph of homogenized and micronized nanomodal structure, b) brightfield TEM micrograph showing detail of matrix grains .
FIG. 9 is an X-ray diffraction pattern of the nanomodal structure of Alloy 1 after hot rolling, a) is experimental data, and b) Rietveld micronization analysis.
10 is a microstructure of Alloy 1 after cold rolling to 1.2 mm thickness, a) is backscatter SEM micrograph showing high strength nanomodal structure after cold rolling, and b) brightfield TEM microscope showing detail of matrix grains It is a photograph.
FIG. 11 is an X-ray diffraction pattern of the high-strength nanomodal structure of Alloy 1 after cold rolling, a) is experimental data, and b) Rietvelt micronization analysis.
FIG. 12 is a brightfield TEM micrograph of the microstructure of Alloy 1 showing the recrystallized modal structure after hot rolling, cold rolling and annealing at 850 ° C. for 5 minutes, a) is a low magnification image, b) is the crystal structure of austenite phase Is a high magnification image with a selected electron diffraction pattern.
FIG. 13 is a backscatter SEM micrograph of the microstructure of Alloy 1 showing the recrystallized modal structure after hot rolling, cold rolling and annealing at 850 ° C. for 5 minutes, where a) is a low magnification image and b) a high magnification image.
FIG. 14 is an X-ray diffraction pattern for the recrystallized modal structure of Alloy 1 after annealing, a) is experimental data and b) Rietvelt refinement analysis.
FIG. 15 is a brightfield TEM micrograph of the microstructure of Alloy 1 showing micronized high strength nanomodal structures (mixed microcomponent structures) formed after tensile strain, a) having large grains and micronized grains of non-morphic structure Transformed "pocket"; b) is the refined structure in the "pocket".
FIG. 16 is a backscatter SEM micrograph of the microstructure of Alloy 1 showing refined high strength nanomodal structure (mixed microcomponent structure), a) a low magnification image, b) a high magnification image.
17 is an X-ray diffraction pattern of the refined high-strength nanomodal structure of Alloy 1 after cold deformation, a) is experimental data, and b) is Rietvelt micronization analysis.
18 is a microstructure of solidified alloy 2 cast to a thickness of 50 mm, a) is a backscatter SEM micrograph showing the dendritic properties of the modal structure in the cast state, and b) a bright field showing the details of the matrix grains TEM micrograph.
19 is an X-ray diffraction pattern of the modal structure of Alloy 2 after coagulation, a) is experimental data, and b) Rietveld micronization analysis.
20 is a microstructure of Alloy 2 after hot rolling to 1.7 mm thickness, a) is a backscatter SEM micrograph of homogenized and refined nanomodal structure, b) a brightfield TEM micrograph showing details of matrix grains .
FIG. 21 is an X-ray diffraction pattern for the nanomodal structure of Alloy 2 after hot rolling, a) is experimental data, and b) Rietveld micronization analysis.
FIG. 22 is a microstructure of Alloy 2 after cold rolling to 1.2 mm thickness, a) is a backscatter SEM micrograph showing high strength nanomodal structure after cold rolling, and b) a brightfield TEM microscope showing detail of matrix grains It is a photograph.
FIG. 23 is an X-ray diffraction pattern of the high-strength nanomodal structure of Alloy 2 after cold rolling, a) is experimental data, and b) is Rietvelt micronization analysis.
FIG. 24 is a brightfield TEM micrograph of the microstructure of Alloy 2 showing the recrystallized modal structure after hot rolling, cold rolling and annealing at 850 ° C. for 5 minutes, a) a low magnification image, b) a crystal structure of austenite phase Is a high magnification image with a selected electron diffraction pattern.
FIG. 25 is a backscatter SEM micrograph of the microstructure of Alloy 2 showing the recrystallized modal structure after hot rolling, cold rolling and annealing at 850 ° C. for 10 minutes, a) is a low magnification image, b) a high magnification image.
FIG. 26 is an X-ray diffraction pattern of the recrystallized modal structure of Alloy 2 after annealing, with a) experimental data and b) Rietvelt micronization analysis.
FIG. 27 is a microstructure of Alloy 2 showing micronized high strength nanomodal structure (mixed microcomponent structure) formed after tensile strain, a) is a brightfield TEM micrograph of a transformed “pocket” with micronized grains, b ) Is a backscatter SEM micrograph of the microstructure.
FIG. 28 X-ray is a diffraction pattern for the refined high strength nanomodal structure of Alloy 2 after cold deformation, a) is experimental data, b) Rietvelt micronization analysis.
29 is a tensile characteristic of Alloy 1 at various stages of a laboratory process.
30 is tensile results for alloy 13 at various stages of the laboratory process.
31 is tensile results for alloy 17 at various stages of the laboratory process.
FIG. 32 is tensile properties in the hot rolled state showing complete property reversibility and after each step of the cold rolling / annealing cycle, a) is an alloy, b) is alloy 2. FIG.
FIG. 33 is a schematic diagram of a flexure test showing schematically a flexure apparatus (International Organization for Standardization, 2005) with two supports and one former.
FIG. 34 is an image of a flexural test sample of alloy 1 tested at 180 °, a) is a photograph of a complete set of samples tested at 180 ° without cracking, b) is an enlarged view of the flexure of the sample tested.
35 a) is the tensile test results of punched and EDM cut specimens from selected alloys showing degradation due to punched edge damage, and b) is the tensile curve of selected alloys for EDM cut specimens.
FIG. 36 is an SEM image of the specimen edge of Alloy 1 after a) EDM cutting and b) punching.
FIG. 37 is an SEM image of the microstructure near the edge of Alloy 1 of a) an EDM cut specimen and b) a punched specimen.
38 is a tensile test result for punched specimens of Alloy 1 before and after annealing showing complete recovery of properties from edge damage by annealing. Data for EDM cut specimens for the same alloy are shown for reference.
FIG. 39 is an exemplary tensile stress-strain curve for punched specimens from alloy 1 annealed and not annealed. FIG.
40 is a tensile stress-strain curve showing the reaction of alloy 1 cold rolled to recovery temperatures ranging from 400 ° C. to 850 ° C., a) is a tensile curve, and b) is a yield strength.
FIG. 41 is a brightfield TEM image of a cold rolled alloy 1 sample showing a highly strained and textured high strength nanomodal structure, where a) is a lower magnification image and b) is a higher magnification image.
FIG. 42 is a brightfield TEM image of an alloy 1 sample annealed at 450 ° C. for 10 minutes, showing a highly strained and textured high strength nanomodal structure with no recrystallization, a) a lower magnification image, b) higher Magnification image.
FIG. 43 is a brightfield TEM image of an alloy 1 sample annealed at 600 ° C. for 10 minutes, showing nanoscale grains showing signs of initiation of recrystallization, a) a lower magnification image, b) a higher magnification image to be.
44 is a brightfield TEM image of an alloy 1 sample annealed at 650 ° C. for 10 minutes, showing larger grains showing a greater degree of recrystallization, a) a lower magnification image, b) a higher magnification image to be.
FIG. 45 is a brightfield TEM image of an alloy 1 sample annealed at 700 ° C. for 10 minutes showing recrystallized grains with some non-morphic regions, electron diffraction shows that the recrystallized grains are austenite, a) The lower magnification image, b) is the higher magnification image.
FIG. 46 is a typical time temperature transformation diagram illustrating the reaction of a steel alloy herein with temperature upon annealing. In the heating curve labeled A, the recovery mechanism is activated. In the heating curve labeled B, the recrystallization mechanism is activated.
47 is tensile properties of punched specimens before and after annealing at different temperatures, a) is alloy 1, b) is alloy 9, and c) is alloy 12. FIG.
48 is a schematic diagram of sample locations for structural analysis.
FIG. 49 is a punched E8 sample of Alloy 1 in the punched state, a) is a low magnification image showing triangular strain regions at the punched edge located on the right side of the photograph. In addition, a magnified area for subsequent micrographs is provided, b) is a higher magnification image showing the deformation region, c) is a higher magnification image showing the recrystallized structure away from the deformation region, and d) It is a higher magnification image showing the deformed structure in the deformation region.
50 is a punched E8 sample of alloy 1 after annealing at 650 ° C. for 10 minutes, a) shows the strain region at the edge punched in the upright direction. In addition, a magnified area for subsequent micrographs is provided, b) is a higher magnification image showing the deformation region, c) is a higher magnification image showing the recrystallized structure away from the deformation region, and d) It is a higher magnification image showing the recovered structure in the deformation region.
FIG. 51 is a punched E8 sample of Alloy 1 after annealing at 700 ° C. for 10 minutes, a) shows the strain region at the edge punched in the upright direction. In addition, a magnified area for subsequent micrographs is provided, b) is a higher magnification image showing the deformation region, c) is a higher magnification image showing the recrystallized structure away from the deformation region, and d) It is a higher magnification image showing the recrystallized structure in the deformation region.
52 is tensile properties for specimens from a) alloy 1, b) alloy 9, c) alloy 12 punched at various speeds.
FIG. 53 is the HER result of Alloy 1 for the punched and milled holes. FIG.
54 is a cut plane for SEM microscopy and microhardness measurement samples from HER tested specimens.
55 is a schematic view of the microhardness measurement position.
FIG. 56 is a microhardness measurement profile of an alloy 1 HER tested sample with a) EDM cut holes and b) punched holes.
FIG. 57 is a microhardness profile for Alloy 1 at various stages of processing and forming showing progress of edge structure transformation during hole punching and expansion.
58 is microhardness data for HER tested samples from Alloy 1 with punched holes and milled holes. The circle represents the position of the TEM sample relative to the hole edge.
FIG. 59 is a brightfield TEM image of the microstructure of Alloy 1 sheet samples prior to HER testing. FIG.
FIG. 60 is a brightfield TEM micrograph of the microstructure of a HER test sample of Alloy 1 with holes punched at about 1.5 mm from the hole edge (HER = 5%), a) is the major non-morphological structure, b ) Is the "pocket" of the partially metamorphic structure.
FIG. 61 is a brightfield TEM micrograph of the microstructure of a HER tested sample of Alloy 1 with milled holes (HER = 73.6%) at about 1.5 mm from the hole edges of the different regions of a) and b).
FIG. 62 is a Focused Ion Beam (FIB) technique used for precise sampling near the edge of a punched hole of Alloy 1 sample, a) is an FIB technique showing the location of a typical sample of a milled TEM sample, b) is an enlarged view of the TEM sample cut at the indicated position from the hole edge.
FIG. 63 is a brightfield TEM micrograph of the microstructure of a sample of Alloy 1 with holes punched at about 10 microns from the hole edge.
64 is a hole expansion ratio measurement of Alloy 1 with and without annealing of punched holes.
65 is a hole expansion ratio measurement of Alloy 9 with or without annealing of punched holes.
FIG. 66 is a hole expansion ratio measurement of Alloy 12 with or without annealing of punched holes. FIG.
FIG. 67 is a hole expansion ratio measurement of Alloy 13 with or without annealing of punched holes. FIG.
FIG. 68 is a hole expansion ratio measurement of Alloy 17 with or without annealing of punched holes. FIG.
69 is tensile performance of Alloy 1 tested at different edge conditions. Note that the tensile performance is reduced compared to tensile samples with edge conditions punched edge conditions, wire EDM cut, punched, and subsequently annealed (for 10 minutes at 850 ° C.).
70 is edge formability measured by the hole expansion rate response of Alloy 1 as a function of edge condition. Note that the hole in the punched state has lower edge formability compared to the hole under conditions of wire EDM cutting, punching and subsequent annealing (for 10 minutes at 850 ° C.).
71 is a punch speed dependency of alloy 1 edge formability as a function of punch speed measured by hole expansion rate. Note that the hole expansion rate consistently increases as the punch speed increases.
FIG. 72 is the punch rate dependency of alloy 9 edge formability as a function of punch rate measured by hole expansion rate. Note that the hole expansion rate gradually increases following the rapid increase in hole expansion rate up to about 25 mm / sec punch speed.
73 is a punch speed dependency of alloy 12 edge formability as a function of punch speed measured by hole expansion rate. Note that the hole expansion rate continues to increase at punch speeds above 100 mm / second following a rapid increase in hole expansion rate up to about 25 mm / sec punch speed.
74 is a punch speed dependency of commercial dual phase 980 steel edge formability measured by hole expansion rate. Note that the hole expansion rate is consistently 21% ± 3% for commercial dual phase 980 steels at all punch rates tested.
75 is a schematic diagram of non-planar punch shapes of 6 ° tapered (left), 7 ° conical (center), and conical plane (right). All dimensions are in millimeters.
FIG. 76 is the effect of punch shape on alloy 1 at 28 mm / sec, 114 mm / sec, and 228 mm / sec punch speeds. Note that for Alloy 1, the influence of the punch shape is attenuated at a punch speed of 228 mm / sec.
77 is the effect of the punch shape on alloy 9 at 28 mm / sec, 114 mm / sec, and 228 mm / sec punch speeds. Note that 7 ° conical punches and conical planar punches produce the highest hole expansion rates.
78 is the effect of punch shape on alloy 12 at 28 mm / sec, 114 mm / sec, and 228 mm / sec punch speeds. Note that the 7 ° conical punch produces the highest hole expansion rate measured for all alloys at a punch speed of 228 mm / sec.
79 is an influence of the punch shape of alloy 1 at a punch speed of 228 mm / sec. Note that all punch shapes produce approximately the same hole expansion rate of about 21%.
80 is a hole punch speed dependency of commercial steel grade edge formability measured by hole expansion rate.
FIG. 81 shows Paul SK [Paul SK, J Mater Eng Perform 2014; using data for commercial steel grades selected from the same paper together with data of Alloy 1 and Alloy 9; 23: 3610.] Predicted the correlation between uniform elongation and hole expansion.
82 is the measured hole expansion rate of a sample of Alloy 1 as a function of hole expansion speed.
83 is the measured hole expansion rate of a sample of alloy 9 as a function of hole expansion speed.
84 is the measured hole expansion rate of a sample of alloy 12 as a function of hole expansion speed.
85 is an image of the microstructure of the sheet of alloy 9, a) is an SEM image of the microstructure, b) is a higher magnification SEM image of the microstructure, c) is an optical image of the etched surface, and d) is etched Is a higher magnification optical image of the surface.
86 is the hole expansion rate measured as a function of hole punching speed and hole expansion speed for the sheet of alloy 9. FIG.
FIG. 87 is the average percent magnetic field volume (Fe%) of HER tested samples using different hole punching rates and hole expansion rates as a function of distance from the hole edge.
88 is a measured hole expansion rate of samples of Alloy 1, Alloy 9, and Alloy 12 as a function of the hole preparation method.
89 shows a low magnification of the cross section near the hole edge of an alloy 1 sample with a) punched holes, b) EDM cut holes, c) milled holes, and d) holes prepared in different ways prior to the expansion of the laser cut holes. SEM image.
90 shows a cross section near the hole edge of an alloy 1 sample with a) punched holes, b) EDM cut holes, c) milled holes, and d) holes prepared in different ways before expansion with high magnification of laser cut holes. High magnification SEM image of.
91 shows near hole edges of an alloy 1 sample with a) punched holes, b) EDM cut holes, c) milled holes, and d) holes prepared in different ways after expansion during HER testing of laser cut holes. Low magnification SEM image of the cross section.
In FIG. 92, SEM images of the sample cross section near the hole edges after HER testing (ie, after expansion to fracture by cracks) show a) punched holes, b) EDM cut holes, c) milled holes, and d ) Is provided at a higher magnification for samples of alloy 1 with holes prepared by different methods of laser cut holes.

Structure and mechanism

Steel alloys of the present disclosure pass through a unique path of structure formation through certain mechanisms as shown in FIGS. 1A and 1B. Initial structure formation begins with melting the alloy, followed by forming the alloy into a cooled, solidified and modal structure (structure # 1, FIG. 1A). Modal structures may contain precipitates containing ferrite grains (α-Fe), martensite, and borides (if boron is present) and / or carbides (if carbon is present), depending on the chemical nature of the particular alloy. Mainly austenite matrix (? -Fe). The grain size of the modal structure depends on the chemistry and solidification conditions of the alloy. For example, thicker cast-in structures (e.g., thicknesses greater than 2.0 mm) result in relatively slower cooling rates (e.g., cooling rates below 250 K / sec) and relatively larger matrix grain sizes. do. Thus, the thickness may preferably range from 2.0 to 500 mm. The modal structure shows austenitic matrix (γ-Fe) with a grain size and / or dendrite length of 2 to 10,000 μm and precipitates of 0.01 to 5.0 μm in laboratory casting. Matrix grain size and precipitate size can be up to 10 times larger in commercial production, depending on the alloy's chemistry, starting casting thickness and specific process parameters. Steel alloys of the present disclosure having a modal structure typically have the following tensile properties, yield strengths from 144 to 514 MPa, maximum tensile strengths in the range from 411 to 907 MPa, and 3.7 to 24.4, depending on the starting thickness size and the chemical properties of the particular alloy. Represents total malleability in%.

Steel alloys of the present disclosure having a modal structure (structure # 1, FIG. 1A) are nanophases by exposing the steel alloy to one or more thermal and stress cycles, ultimately leading to the formation of nanomodal structure (structure # 2, FIG. 1A). Homogenization and refinement can be achieved through micronization (mechanism # 1, FIG. 1A). More specifically, the modal structure, when formed to a thickness of at least 2.0 mm or formed at a cooling rate of 250 K / sec or less, is preferably heated to a temperature below the solidus temperature (T _m ) to a temperature of 700 ° C. , Strain rate of thickness reduction of 10 ⁻⁶ to 10 ⁴ . Transformation to structure # 2 results in an intermediate homogenized modal structure (structure # 1a, as the steel alloy undergoes mechanical deformation during continuous application of temperature and stress, a thickness reduction that can be configured to occur during hot rolling, and the like. Generated continuously through FIG. 1A).

The nanomodal structure (structure # 2, FIG. 1A) has a primary austenitic matrix (γ-Fe) and, depending on chemistry, additionally ferrite grains (α-Fe) and / or precipitates (eg, borides (boron) Is present) and / or carbides (if carbon is present) Depending on the starting grain size, nanomodal structures are typically primary austenitic based having grain sizes of 1.0 to 100 μm in laboratory casting. Matrix (γ-Fe) and / or 1.0 to 200 nm sized precipitates The matrix grain size and precipitate size can be up to 5 times larger in commercial production depending on the alloy's chemical properties, starting casting thickness and specific process parameters Steel alloys of this disclosure having nanomodal structures typically have the following tensile properties, yield strengths of 264 to 574 MPa, maximum tensile strengths in the range of 921 to 1413 MPa, and 12.0 to 7 7.7% total malleability Structure # 2 is preferably formed with a thickness of 1 mm to 500 mm.

When a steel alloy having a nanomodal structure (structure # 2, FIG. 1A) is stressed near ambient / ambient temperature (eg, 25 ° C. ± 5 ° C.), a dynamic nanophase strengthening mechanism (mechanism # 2, FIG. 1a) is activated leading to the formation of a high strength nanomodal structure (structure # 3, FIG. 1a). Preferably, the stress is at a level higher than each yield strength of the alloy in the range of 250 to 600 MPa, depending on the chemical nature of the alloy. High-strength nanomodal structures typically exhibit a ferritic matrix (α-Fe), which, depending on the alloy's chemical properties, additionally has austenite grains (γ-Fe), and borides (if boron is present) and / or carbides ( May contain precipitate grains, which may include carbon). Note that the strengthening transformation occurs during deformation under applied stress, which defines mechanism # 2 as a dynamic process in which the metastable austenitic phase (γ-Fe) is transformed into ferrite (α-Fe) containing precipitates. Note that, depending on the starting chemistry, some of the austenite is stable and will not be transformed. Typically, 5 to 95 volume percent of the matrix is transformed. High-strength nanomodal structures typically show ferrite matrix (α-Fe) with matrix grain sizes of 25 nm to 50 μm and precipitate grains of 1.0 to 200 nm in laboratory casting. Matrix grain size and precipitate size can be up to 2 times larger in commercial production, depending on the alloy's chemistry, starting casting thickness and specific process parameters. Steel alloys of this disclosure having high strength nanomodal structures typically exhibit the following tensile properties, yield strength of 718 to 1645 MPa, maximum tensile strength in the range of 1356 to 1831 MPa, and total conductivity of 1.6 to 32.8%. Structure # 3 is preferably formed to a thickness of 0.2 to 25.0 mm.

The high-strength nanomodal structure (structures # 3, Figures 1A and 1B), when heated below the melting point of the alloy, transforms ferrite grains into austenite to form a recrystallized modal structure (structure # 4, Figure 1B). Has the ability to undergo recrystallization (mechanism # 3, FIG. 1B). Partial solubility of nanoscale precipitates also occurs. Depending on the chemical nature of the alloy, borides and / or carbides may be present in the material. The preferred temperature range for full transformation is from 650 ° C. to the T _m of the particular alloy. When recrystallized, structure # 4 contains little dislocations or twins, and stacking defects may be found in some recrystallized grains. Note that recovery mechanisms may occur at lower temperatures of 400-650 ° C. The recrystallized modal structure (structure # 4, FIG. 1B) typically shows primary austenitic matrix (γ-Fe) having a grain size of 0.5 to 50 μm and precipitate grains of 1.0 to 200 nm in laboratory casting. Matrix grain size and precipitate size can be up to 2 times larger in commercial production, depending on the alloy's chemistry, starting casting thickness and specific process parameters. Steel alloys of this disclosure having a recrystallized modal structure typically exhibit the following tensile properties, yield strengths from 197 to 1372 MPa, maximum tensile strengths in the range from 799 to 1683 MPa, and total conductivity of 10.6 to 86.7%.

Steel alloys of the present disclosure having a recrystallized modal structure (structure # 4, FIG. 1B) are characterized by near ambient / ambient temperatures (e.g., causing the formation of refined high strength nanomodal structure (structure # 5, FIG. 1B). Nanophase refinement and reinforcement (mechanism # 4, FIG. 1B) when subjected to stresses above yield at 25 ° C. ± 5 ° C.). Preferably, the stress for initiating mechanism # 4 is at a level above the yield strength in the range of 197 to 1372 MPa. Similar to mechanism # 2, the nanophase refinement and strengthening (mechanism # 4, FIG. 1B) is a metastable austenitic phase that transforms into ferrite containing precipitates and is generally further grain compared to structure # 3 for the same alloy. It is a dynamic process that causes micronization. One of the features of the refined high-strength nanomodal structure (structure # 5, FIG. 1B) is that significant refinement occurs in the "pockets" of the irregularly distributed microstructures while other regions remain untransformed during phase transformation. Note that, depending on the starting chemistry, some of the austenite is stable and the regions containing stabilized austenite will not be transformed. Typically, 5 to 95 volume percent of the matrix is transformed within the distributed "pocket". Depending on the chemistry of the alloy, borides (if boron is present) and / or carbides (if carbon is present) may be present in the material. The non-transformation part of the microstructure is represented by austenite-based grains (γ-Fe) of 0.5 to 50 μm in size and may also contain distributed precipitates of to 200 nm in size. These highly strained austenitic grains contain a relatively large number of dislocations by existing dislocation processes occurring during deformation, thereby increasing the ratio of dislocations (10 ⁸ to 10 ¹⁰ mm ⁻² ). The transformed portion of the microstructure during deformation is represented by refined ferrite grains (α-Fe) with additional precipitates through nanophase refinement and reinforcement (mechanism # 4, FIG. 1B). In laboratory casting, the size of the micronized grains of ferrite (α-Fe) is in the range of 50 to 2000 nm, and the size of the precipitate is 1 to 200 nm. Matrix grain size and precipitate size can be up to 2 times larger in commercial production, depending on the alloy's chemistry, starting casting thickness and specific process parameters. The size of the "pockets" of metamorphized and highly microstructured microstructures typically ranges from 0.5 to 20 μm. The volume fraction of the transformed regions relative to the non-transformed regions in the microstructure can be varied, typically at a ratio of 95: 5 to 5:95, respectively, by varying the chemical properties of the alloy including austenite stability. Steel alloys of the present disclosure having a refined high strength nanomodal structure typically exhibit the following tensile properties, yield strength of 718 to 1645 MPa, maximum tensile strength in the range of 1356 to 1831 MPa, and total conductivity of 1.6 to 32.8%.

The steel alloys of the present disclosure, which have then refined high strength nanomodal structures (structure # 5, FIG. 1B), may be exposed to elevated temperatures to return to the formation of recrystallized modal structures (structure # 4, FIG. 1B). Typical temperature ranges for complete transformation are 650 ° C. to the T _m of the particular alloy (illustrated in FIG. 1B), and lower temperatures from 400 ° C. to less than 650 ° C. may activate the recovery mechanism and cause partial recrystallization. The stress portion and heating are repeated a plurality of times to obtain the desired product shape, including, but not limited to, relatively thin sheets having target properties, relatively small diameter tubes or rods, complex shaped end parts, and the like. Can be. Thus the final thickness of the material may range from 0.2 to 25 mm. Note that cubic precipitates may be present in the steel alloys of the present disclosure having a Fm3m (# 225) space group at all stages. _Bihexagonal pyramidal class hexagonal phase and / or hexagonal P6bar2C with P6 _3mc spatial group (# 186) due to _{dynamic nanophase} enhancement mechanisms (mechanism # 2) and / or deformation through _nanophase refinement and enhancement (mechanism # 4) Additional nanoscale precipitates may be formed which are represented by the bi-three-way bipyramid class with space group (# 190). Precipitate properties and volume fractions depend on alloy composition and process history. The size of the nanoprecipitates may range from 1 nm to several tens of nanometers but in most cases is less than 20 nm. The volume fraction of the precipitate is generally less than 20%.

Mechanism during sheet production through slab casting

The structures and effective mechanisms for the steel alloys herein are applicable to commercial production using existing process flows. See FIG. 2. Steel slabs are usually manufactured by continuous casting with a number of subsequent process variations to obtain the shape of the final product, which is the coil of the sheet. FIG. 2 shows the detailed structural evolution of the steel alloy herein from casting to the final product for each stage of processing from the slab to the sheet product.

Formation of a modal structure (structure # 1) in the steel alloy herein occurs during alloy solidification. The modal structure preferably heats the alloy of the specification to a temperature in the range above its melting point and in the range of 1100 ° C. to 2000 ° C., preferably corresponding to cooling in the range of ¹ × 10 ³ to 1 × 10 ⁻³ K / sec. It can be formed by cooling below the melting temperature of. The thickness of the casting depends on the production method, where thin slab castings are typically 20 to 150 mm thick and thick slab castings are typically 150 to 500 mm thick. Thus, the casting thickness can range from 20 to 500 mm, and all values within the range can increase by 1 mm. Thus, the casting thicknesses are 21 mm, 22 mm, 23 mm and the like, up to 500 mm.

Hot rolling of slabs solidified from the alloy is the next process step to produce transfer bars in the case of thick slab casting or coils in the case of thin slab casting. During this process, the modal structure is continuously transformed into a partially homogenized and then fully homogenized modal structure (structure # 1a) through nanophase refinement (mechanism # 1). Upon completion of homogenization and thereby miniaturization, a nanomodal structure (structure # 2) is formed. The resulting hot band coil, which is the product of a hot rolling process, is typically in the range of 1 to 20 mm in thickness.

Cold rolling is a sheet production method that is widely used to achieve target thicknesses for specific applications. For AHSS, thinner specifications in the 0.4 to 2 mm range are generally targeted. To achieve a thinner specification thickness, cold rolling may be applied through multiple passes, with or without intermediate annealing between the passes. Typical reductions per pass are from 5 to 70% depending on material properties and equipment capacity. The number of passes before intermediate annealing also depends on the material properties and the level of strain hardening during cold deformation. In the case of the steel alloys herein, cold rolling causes dynamic nanophase reinforcement (mechanism # 2) and causes extensive strain hardening of the resulting sheet and the formation of high strength nanomodal structures (structure # 3). The properties of the cold rolled sheet from the alloys herein depend on the chemistry of the alloy and are controlled by the reduction of cold rolling to produce a completely cold rolled (ie hard) product, or by various properties (ie, ¼, ½, ¾ hard, etc.). Depending on the specific process flow, especially the starting thickness and the amount of reduction in hot rolling specifications, annealing is often necessary to restore the malleability of the material to reduce additional cold rolling specifications. The intermediate coil can be annealed using conventional methods such as batch anneal or continuous anneal lines. Cold strained high strength nanomodal structures (structure # 3) for the steel alloys herein are recrystallized (mechanism # 3) during annealing, leading to the formation of recrystallized modal structures (structure # 4). At this stage, the recrystallized coil may be the final product with advanced combinations of properties depending on the alloy's chemistry and target market. If a much thinner sheet is required, the recrystallized coil can be subjected to additional cold rolling to achieve a target thickness that can be realized by one or more cold rolling / annealing cycles. Further cold deformation from the alloys of the present disclosure with a recrystallized modal structure (structure # 4) leads to structural transformation into nanostructured high strength nanomodal structure (structure # 5) through nanophase refinement and reinforcement (mechanism # 4). do. As a result, a completely rigid coil with a final specification and a refined high strength nanomodal structure (structure # 5) can be formed, or in the case of annealing as the final stage of the cycle, the final specification and recrystallized modal structure (structure # Coils of sheets having 4) may be manufactured. Micronized high strength nanomodal structures (structure # 5) when the coils of sheets recrystallized from the alloys of the present disclosure are used for the production of finished parts by any type of cold deformation, such as cold stamping, hydroforming, roll forming, etc. Will be present in the final product / part. The final product can exist in many different forms, including sheets, plates, strips, pipes, and tubes, and numerous complex parts can be fabricated through a variety of metalworking processes.

Edge Formability Mechanism

The circulating nature of this phase transformation from the recrystallized modal structure (structure # 4) to the refined high-strength nanomodal structure (structure # 5) and then back to the recrystallized modal structure (structure # 4) is described herein. One of the unique phenomena and features. As mentioned above, this cycling characteristic is applicable during commercial production of the sheet, especially in the case of AHSS which requires a thinner specification thickness (eg thickness in the range of 0.2 to 25 mm). This reversible mechanism is also applicable for a wide range of industrial applications of the steel alloys herein. In the case of the steel alloys herein, an excellent combination of bulk sheet formability is demonstrated, as evidenced by the tensile and bending properties in these applications, but the unique cycling properties of phase transformation can be a significant limiting factor for other AHSS. Enable edge formability. Table 1 below summarizes the structural and performance characteristics through stressing and heating cycles available through nanophase refinement and reinforcement (mechanism # 4). Methods of using these structures and mechanisms to produce excellent combinations of bulk sheet and edge formability will be described later.

Table 1. Structure and performance with stress addition / heating cycle

Attribute / mechanism Structure # 4
Recrystallized Modal Structure Structure # 5
Micronized high strength nanomodal structure Kinky Perverted "Pocket" rescue
formation Recrystallization at elevated temperatures in cold worked materials Retained austenitic grains Nanophase refinement and reinforcement mechanisms generated by the application of mechanical stresses in the "pockets" of distributed microstructures transformation Recrystallization of cold strained iron matrix Random precipitation Austenitic transformation into stresses caused by ferrites and precipitates Effective phase Austenitic, optionally ferrite, precipitates Austenitic, optionally precipitates Ferrite, optionally austenite, precipitates Matrix grain size 0.5 to 50 μm 0.5 to 50 μm 50 to 2000 nm Precipitate size 1 to 200 nm 1 to 200 nm 1 to 200 nm Tensile reaction Actual with properties achieved based on the formation of structures and the fraction of transformation Actual with properties achieved based on the formation of structures and the fraction of transformation Yield strength 197 to 1372 MPa 718 to 1645 MPa Tensile strength 799 to 1683 MPa 1356 to 1831 MPa Total elongation 6.6 to 86.7% 1.6 to 32.8%

main text

The chemical compositions of the alloys herein are shown in Table 2 which provides the preferred atomic ratios used.

Table 2 Chemical Composition of Alloys

alloy Fe Cr Ni Mn Cu B Si C Alloy 1 75.75 2.63 1.19 13.86 0.65 0.00 5.13 0.79 Alloy 2 73.99 2.63 1.19 13.18 1.55 1.54 5.13 0.79 Alloy 3 77.03 2.63 3.79 9.98 0.65 0.00 5.13 0.79 Alloy 4 78.03 2.63 5.79 6.98 0.65 0.00 5.13 0.79 Alloy 5 79.03 2.63 7.79 3.98 0.65 0.00 5.13 0.79 Alloy 6 78.53 2.63 3.79 8.48 0.65 0.00 5.13 0.79 Alloy 7 79.53 2.63 5.79 5.48 0.65 0.00 5.13 0.79 Alloy 8 80.53 2.63 7.79 2.48 0.65 0.00 5.13 0.79 Alloy 9 74.75 2.63 1.19 14.86 0.65 0.00 5.13 0.79 Alloy 10 75.25 2.63 1.69 13.86 0.65 0.00 5.13 0.79 Alloy 11 74.25 2.63 1.69 14.86 0.65 0.00 5.13 0.79 Alloy 12 73.75 2.63 1.19 15.86 0.65 0.00 5.13 0.79 Alloy 13 77.75 2.63 1.19 11.86 0.65 0.00 5.13 0.79 Alloy 14 74.75 2.63 2.19 13.86 0.65 0.00 5.13 0.79 Alloy 15 73.75 2.63 3.19 13.86 0.65 0.00 5.13 0.79 Alloy 16 74.11 2.63 2.19 13.86 1.29 0.00 5.13 0.79 Alloy 17 72.11 2.63 2.19 15.86 1.29 0.00 5.13 0.79 Alloy 18 78.25 2.63 0.69 11.86 0.65 0.00 5.13 0.79 Alloy 19 74.25 2.63 1.19 14.86 1.15 0.00 5.13 0.79 Alloy 20 74.82 2.63 1.50 14.17 0.96 0.00 5.13 0.79 Alloy 21 75.75 1.63 1.19 14.86 0.65 0.00 5.13 0.79 Alloy 22 77.75 2.63 1.19 13.86 0.65 0.00 3.13 0.79 Alloy 23 76.54 2.63 1.19 13.86 0.65 0.00 5.13 0.00 Alloy 24 67.36 10.70 1.25 10.56 1.00 5.00 4.13 0.00 Alloy 25 71.92 5.45 2.10 8.92 1.50 6.09 4.02 0.00 Alloy 26 61.30 18.90 6.80 0.90 0.00 5.50 6.60 0.00 Alloy 27 71.62 4.95 4.10 6.55 2.00 3.76 7.02 0.00 Alloy 28 62.88 16.00 3.19 11.36 0.65 0.00 5.13 0.79 Alloy 29 72.50 2.63 0.00 15.86 1.55 1.54 5.13 0.79 Alloy 30 80.19 0.00 0.95 13.28 1.66 2.25 0.88 0.79 Alloy 31 77.65 0.67 0.08 13.09 1.09 0.97 2.73 3.72 Alloy 32 78.54 2.63 1.19 13.86 0.65 0.00 3.13 0.00 Alloy 33 83.14 1.63 8.68 0.00 1.00 4.76 0.00 0.79 Alloy 34 75.30 2.63 1.34 14.01 0.80 0.00 5.13 0.79 Alloy 35 74.85 2.63 1.49 14.16 0.95 0.00 5.13 0.79

As can be seen from the above, the alloy of the present specification is an iron-based metal alloy having 50 atomic% or more of Fe. More preferably, the alloy herein may be described as including, consisting essentially of, or consisting of the following elements, expressed in atomic percent: Fe (61.30 to 83.14 atomic percent) ); Si (0 to 7.02 atomic%); Mn (0-15.86 atomic%); B (0 to 6.09 atomic%); Cr (0-18.90 atomic%); Ni (0 to 8.68 atomic%); Cu (0-2.00 atomic%); C (0-3.72 atomic%). In addition, the alloy of the present specification may be understood to include Fe and 4 or more, or 5 or more, or 6 or more elements selected from Si, Mn, B, Cr, Ni, Cu or C. Most preferably, the alloys herein comprise Fe, consist essentially of, or consist of Fe at a level of at least 50 atomic% with Si, Mn, B, Cr, Ni, Cu and C.

Alloy laboratory process

The laboratory process of the alloys of Table 2 was performed to model each stage of industrial production on a much smaller scale. The main steps in this process include: casting, tunnel furnace heating, hot rolling, cold rolling, and annealing.

casting

The alloys were weighed into charges in the range of 3,000 to 3,400 grams using commercially available ferroditive powders with known chemical properties and impurity contents according to the atomic ratios of Table 2. The charge was loaded into a zirconia coated silica crucible installed in an Indutherm VTC800V vacuum tilt casting machine. The casting machine was then evacuated to cast and melt chambers and backfilled with argon several times before casting to prevent oxidation of the melt. The melt was heated with a 14 kHz RF induction coil until it melted completely for about 5.25 to 6.5 minutes, depending on the alloy composition and the mass of the charge. After melting of the last solid was observed, it was further heated for 30 to 45 seconds to provide overheating to ensure homogenization of the melt. The casting machine then emptied the melting and casting chamber and tilted the crucible to pour the melt into a 50 mm thick, 75 to 80 mm wide, and 125 mm deep channel of the water cooled copper die. The melt was cooled under vacuum for 200 seconds and then the chamber was filled with argon to atmospheric pressure. A photograph of an embodiment of a lab cast slab from two different alloys is shown in FIG. 3.

Tunnel furnace heating

The laboratory slab was charged into a Lucifer EHS3GT-B18 furnace for heating before hot rolling. The set point of the furnace is in the range of 1100 ° C. to 1250 ° C., depending on the alloy melting point. This was soaked for 40 minutes before hot rolling to ensure that the slab reached the target temperature. Between the hot rolling and the hot rolling the slab was returned to the furnace and reheated for 4 minutes.

Hot rolling

The preheated slab was pushed from the tunnel furnace into the Fenn Model 061 2 high pressure smoke. The 50 mm slab is preferably air cooled after hot rolling through 5 to 8 passes through high pressure steam. After the first pass the angular slab was reduced by 80 to 85% to a final thickness of 7.5 to 10 mm. Each obtained sheet was cut after cooling and hot rolled through 190 mm of the bottom through 3 to 4 additional passes through the rolling mill to further shrink to a final thickness of 1.6 to 2.1 mm between 72 and 84%. A photograph of an example of a lab cast slab from two different alloys after hot rolling is shown in FIG. 4.

Cold rolled

The sheet obtained after hot rolling was media blasted with aluminum oxide to remove the rolling scale, which was then cold rolled on a Fenn Model 061 2 high pressure steamer. Cold rolling takes several passes to reduce the thickness of the sheet, typically to a target thickness of 1.2 mm. The hot rolled sheet was fed into the rolling mill continuously decreasing the roll gap until the minimum gap was reached. If the material did not reach the specification target, additional passes were used with minimum clearance until the 1.2 mm thickness was reached. Many passes were applied due to the limitations of laboratory mill capacity. An example photograph of a cold rolled sheet from two different alloys is shown in FIG. 5.

Annealing

After cold rolling, the tensile specimens were cut from the cold rolled sheet via wire electrical discharge machining (EDM). These specimens were then annealed with the different parameters reported in Table 3. Annealing (1a, 1b, 2b) was performed in a Lucifer 7HT-K12 box furnace. Annealing (2a and 3) was performed in a Camco Model G-ATM-12FL furnace. Air normalized specimens were removed from the furnace at the end of the cycle and cooled to room temperature in air. For furnace cooled specimens, the furnace was de-energized at the end of the annealing to allow the sample to cool with the furnace. Note that the heat treatment was chosen for the demonstration and is not intended to limit the scope. High temperature treatment up to the melting point of each alloy is possible.

[Table 3] Annealing Parameters

Annealing heating Temperature Residence time Cooling atmosphere 1a Preheated furnace 850 ℃ 5 minutes Air Normalized Air + argon 1b Preheated furnace 850 ℃ 10 minutes Air Normalized Air + argon 2a 20 ℃ / hour 850 ℃ 360 minutes 45 ℃ / hour up to 500 ℃
Furnace Cooling Next Hydrogen + argon 2b 20 ℃ / hour 850 ℃ 360 minutes 45 ℃ / hour up to 500 ℃
Air normalized to Air + argon 3 20 ℃ / hour 1200 ℃ 120 minutes Furnace cooling Hydrogen + argon

Alloy properties

Thermal analysis of the alloys of the present disclosure was performed on a solidified cast slab using a Netzsch Pegasus 404 Differential Scanning Calorimeter (DSC). An alloy sample was charged into an alumina crucible and charged into the DSC. This DSC then emptied the chamber and backfilled with argon to atmospheric pressure. A constant purge of argon was then initiated and a zirconium getter was installed in the gas flow path to further reduce the amount of oxygen in the system. The sample was heated to complete melting, cooled to complete solidification, and then reheated to 10 ° C./min through melting. The solidus, liquidus, and peak temperatures were measured from the second melt to ensure a representative measurement of the material in equilibrium. In the alloys listed in Table 2, melting occurs in one or more stages with an initial melting of about 1111 ° C. and a final melting temperature of about 1476 ° C. (Table 4), depending on the alloy's chemistry. The change in melt behavior reflects the complex phase formation upon solidification of the alloy, depending on the alloy's chemistry.

Table 4 Differential thermal analysis data for melt behavior

alloy Solidus temperature
(℃) Liquidus temperature
(℃) Melt Peak # 1 (° C) Melt Peak # 2 (° C) Melt Peak # 3 (° C) Alloy 1 1390 1448 1439 Alloy 2 1157 1410 1177 1401 Alloy 3 1411 1454 1451 Alloy 4 1400 1460 1455 Alloy 5 1415 1467 1464 Alloy 6 1416 1462 1458 Alloy 7 1421 1467 1464 Alloy 8 1417 1469 1467 Alloy 9 1385 1446 1441 Alloy 10 1383 1442 1437 Alloy 11 1384 1445 1442 Alloy 12 1385 1443 1435 Alloy 13 1401 1459 1451 Alloy 14 1385 1445 1442 Alloy 15 1386 1448 1441 Alloy 16 1384 1439 1435 Alloy 17 1376 1442 1435 Alloy 18 1395 1456 1431 1449 1453 Alloy 19 1385 1437 1432 Alloy 20 1374 1439 1436 Alloy 21 1391 1442 1438 Alloy 22 1408 1461 1458 Alloy 23 1403 1452 1434 1448 Alloy 24 1219 1349 1246 1314 1336 Alloy 25 1186 1335 1212 1319 Alloy 26 1246 1327 1268 1317 Alloy 27 1179 1355 1202 1344 Alloy 28 1158 1402 1176 1396 Alloy 29 1159 1448 1168 1439 Alloy 30 1111 1403 1120 1397 Alloy 31 1436 1475 1464 Alloy 32 1436 1476 1464 Alloy 33 1153 1418 1178 1411 Alloy 34 1397 1448 1445 Alloy 35 1394 1444 1441

The density of the alloy was measured in a 9 mm thick section of hot rolled material using the Archimedes method with a balance specifically configured to allow weighing in both air and distilled water. The density of each alloy is shown in Table 5, it can be seen that the range of 7.57 to 7.89 g / cm ³ . The accuracy of this technique is ± 0.01 g / cm ³ .

[Table 5] Density of alloy

alloy density
(g / cm ³ ) alloy density
(g / cm ³ ) Alloy 1 7.78 Alloy 19 7.77 Alloy 2 7.74 Alloy 20 7.78 Alloy 3 7.82 Alloy 21 7.78 Alloy 4 7.84 Alloy 22 7.87 Alloy 5 7.76 Alloy 23 7.81 Alloy 6 7.83 Alloy 24 7.67 Alloy 7 7.79 Alloy 25 7.71 Alloy 8 7.71 Alloy 26 7.57 Alloy 9 7.77 Alloy 27 7.67 Alloy 10 7.78 Alloy 28 7.73 Alloy 11 7.77 Alloy 29 7.89 Alloy 12 7.77 Alloy 30 7.78 Alloy 13 7.80 Alloy 31 7.89 Alloy 14 7.78 Alloy 32 7.89 Alloy 15 7.79 Alloy 33 7.78 Alloy 16 7.79 Alloy 34 7.77 Alloy 17 7.77 Alloy 35 7.78 Alloy 18 7.79

Tensile properties were measured on an Instron 3369 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Tensile properties of the alloys listed in Table 2 after annealing with the parameters listed in Table 3 are shown in Tables 6-10 below. The maximum tensile strength value may range from 799 to 1683 MPa at a tensile elongation of 6.6 to 86.7%. Yield strength ranges from 197 to 978 MPa. The mechanical property values of the steel alloys herein will depend on the chemical properties and process conditions of the alloy. Changes in heat treatment also show possible property changes through treatment of the chemical properties of certain alloys.

Table 6 Tensile data of selected alloys after heat treatment 1a

alloy Yield strength
(MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 443 1212 51.1 458 1231 57.9 422 1200 51.9 Alloy 2 484 1278 48.3 485 1264 45.5 479 1261 48.7 Alloy 3 458 1359 43.9 428 1358 43.7 462 1373 44.0 Alloy 4 367 1389 36.4 374 1403 39.1 364 1396 32.1 Alloy 5 510 1550 16.5 786 1547 18.1 555 1552 16.2 Alloy 6 418 1486 34.3 419 1475 35.2 430 1490 37.3 Alloy 7 468 1548 20.2 481 1567 20.3 482 1545 19.3 Alloy 8 851 1664 13.6 848 1683 14.0 859 1652 12.9 Alloy 9 490 1184 58.0 496 1166 59.1 493 1144 56.6 Alloy 10 472 1216 60.5 481 1242 58.7 470 1203 55.9 Alloy 11 496 1158 65.7 498 1155 58.2 509 1154 68.3 Alloy 12 504 1084 48.3 515 1105 70.8 518 1106 66.9 Alloy 13 478 1440 41.4 486 1441 40.7 455 1424 42.0 Alloy 22 455 1239 48.1 466 1227 55.4 460 1237 57.9 Alloy 23 419 1019 48.4 434 1071 48.7 439 1084 47.5 Alloy 28 583 932 61.5 594 937 60.8 577 930 61.0 Alloy 29 481 1116 60.0 481 1132 55.4 486 1122 56.8 Alloy 30 349 1271 42.7 346 1240 36.2 340 1246 42.6 Alloy 31 467 1003 36.0 473 996 29.9 459 988 29.5 Alloy 32 402 1087 44.2 409 1061 46.1 420 1101 44.1

Table 7 Tensile data of selected alloys after heat treatment 1b

alloy Yield strength (MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 487 1239 57.5 466 1269 52.5 488 1260 55.8 Alloy 2 438 1232 49.7 431 1228 49.8 431 1231 49.4 Alloy 9 522 1172 62.6 466 1170 61.9 462 1168 61.3 Alloy 12 471 1115 63.3 458 1102 69.3 454 1118 69.1 Alloy 13 452 1408 40.5 435 1416 42.5 432 1396 46.0 Alloy 14 448 1132 64.4 443 1151 60.7 436 1180 54.3 Alloy 15 444 1077 66.9 438 1072 65.3 423 1075 70.5 Alloy 16 433 1084 67.5 432 1072 66.8 423 1071 67.8 Alloy 17 420 946 74.6 421 939 77.0 425 961 74.9 Alloy 19 496 1124 67.4 434 1118 64.8 435 1117 67.4 Alloy 20 434 1154 58.3 457 1188 54.9 448 1187 60.5 Alloy 21 421 1201 54.3 427 1185 59.9 431 1191 47.8 Alloy 24 554 1151 23.5 538 1142 24.3 562 1151 24.3 Alloy 25 500 1274 16.0 502 1271 15.8 483 1280 16.3 Alloy 26 697 1215 20.6 723 1187 21.3 719 1197 21.5 Alloy 27 538 1385 20.6 574 1397 20.9 544 1388 21.8 Alloy 33 978 1592 6.6 896 1596 7.2 953 1619 7.5 Alloy 34 467 1227 56.7 476 1232 52.7 462 1217 51.6 Alloy 35 439 1166 56.3 438 1166 59.0 440 1177 58.3

Table 8 Tensile data of selected alloys after heat treatment 2a

alloy
Yield strength
(MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 2 367 1174 46.2 369 1193 45.1 367 1179 50.2 Alloy 30 391 1118 55.7 389 1116 60.5 401 1113 59.5 Alloy 32 413 878 17.6 399 925 20.5 384 962 21.0 Alloy 31 301 1133 37.4 281 1125 38.7 287 1122 39.0

Table 9 Tensile data of selected alloys after heat treatment 2b

alloy Yield strength (MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 396 1093 31.2 383 1070 30.4 393 1145 34.7 Alloy 2 378 1233 49.4 381 1227 48.3 366 1242 47.7 Alloy 3 388 1371 41.3 389 1388 42.6 Alloy 4 335 1338 21.7 342 1432 30.1 342 1150 17.3 Alloy 5 568 1593 15.2 595 1596 13.1 735 1605 14.6 Alloy 6 399 1283 17.5 355 1483 24.8 386 1471 23.8 Alloy 7 605 1622 16.3 639 1586 15.2 Alloy 8 595 1585 13.6 743 1623 14.1 791 1554 13.9 Alloy 9 381 1125 53.3 430 1111 44.8 369 1144 51.1 Alloy 10 362 1104 37.8 369 1156 43.5 Alloy 11 397 1103 52.4 390 1086 50.9 402 1115 50.4 Alloy 12 358 1055 64.7 360 1067 64.4 354 1060 62.9 Alloy 13 362 982 17.3 368 961 16.3 370 989 17.0 Alloy 14 385 1165 59.0 396 1156 55.5 437 1155 57.9 Alloy 15 357 1056 70.3 354 1046 68.2 358 1060 70.7 Alloy 16 375 1094 67.6 384 1080 63.4 326 1054 65.2 Alloy 17 368 960 77.2 370 955 77.9 358 951 75.9 Alloy 18 326 1136 17.3 338 1192 19.1 327 1202 18.5 Alloy 19 386 1134 64.5 378 1100 60.5 438 1093 52.5 Alloy 20 386 1172 56.2 392 1129 42.0 397 1186 57.8 Alloy 21 363 1141 49.0 Alloy 22 335 1191 45.7 322 1189 41.5 348 1168 34.5 Alloy 23 398 1077 44.3 367 1068 44.8 Alloy 24 476 1149 28.0 482 1154 25.9 495 1145 26.2 Alloy 25 452 1299 16.0 454 1287 15.8 441 1278 15.1 Alloy 26 619 1196 26.6 615 1189 26.2 647 1193 26.1 Alloy 27 459 1417 17.3 461 1410 16.8 457 1410 17.1 Alloy 28 507 879 52.3 498 874 42.5 493 880 44.7 Alloy 32 256 1035 42.3 257 1004 42.1 257 1049 34.8 Alloy 33 830 1494 8.4 862 1521 8.1 877 1519 8.8 Alloy 34 388 1178 59.8 384 1197 57.7 370 1177 59.1 Alloy 35 367 1167 58.5 369 1167 58.4 375 1161 59.7

Table 10 Tensile data of selected alloys after heat treatment 3

alloy Yield strength (MPa) Tensile strength (MPa) Tensile Elongation (%) Alloy 1 238 1142 47.6 233 1117 46.3 239 1145 53.0 Alloy 3 266 1338 38.5 Applicable 1301 37.7 Applicable 1291 35.6 Alloy 4 Applicable 1353 27.7 Applicable 1337 26.1 Applicable 1369 29.0 Alloy 5 511 1462 12.5 558 1399 10.6 Alloy 6 311 1465 24.6 308 1467 21.8 308 1460 25.0 Alloy 7 727 1502 12.5 639 1474 11.3 685 1520 12.4 Alloy 8 700 1384 12.3 750 1431 13.3 Alloy 9 234 1087 55.0 240 1070 56.4 242 1049 58.3 Alloy 10 229 1073 50.6 228 1082 56.5 229 1077 54.2 Alloy 11 232 1038 63.8 232 1009 62.4 228 999 66.1 Alloy 12 229 979 65.6 228 992 57.5 222 963 66.2 Alloy 13 277 1338 37.3 261 1352 35.9 272 1353 34.9 Alloy 14 228 1074 58.5 239 1077 54.1 230 1068 49.1 Alloy 15 206 991 60.9 208 1024 58.9 Alloy 16 199 1006 57.7 242 987 53.4 208 995 57.0 Alloy 17 222 844 72.6 197 867 64.9 213 869 66.5 Alloy 18 288 1415 32.6 300 1415 32.1 297 1421 29.6 Alloy 19 225 1032 58.5 213 1019 61.1 214 1017 58.4 Alloy 20 233 1111 57.3 227 1071 53.0 230 1091 49.4 Alloy 21 238 1073 50.6 228 1069 56.5 246 1110 52.0 Alloy 22 217 1157 47.0 236 1154 46.8 218 1154 47.7 Alloy 23 208 979 45.4 204 984 43.4 204 972 38.9 Alloy 28 277 811 86.7 279 802 86.0 277 799 82.0 Alloy 32 203 958 33.3 206 966 39.5 210 979 36.3 Alloy 34 216 1109 52.8 230 1144 55.9 231 1123 52.3 Alloy 35 230 1104 51.7 231 1087 59.0 220 1084 54.4

Example

Example # 1: Structural Development Pathway of Alloy 1

A 50 mm thick laboratory slab was cast from alloy 1 as described in the text of this application, which was hot rolled, cold rolled and annealed at 850 ° C. for 5 minutes by laboratory treatment. The microstructure of the alloy at each process step was examined by SEM, TEM and X-ray analysis.

For SEM irradiation, the cross section of the slab sample was polished with a small grit SiC abrasive paper and then polished gradually with diamond media paste up to 1 μm. Final polishing was performed using a 0.02 μm grit SiO ₂ solution. Microstructure was examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. To prepare TEM specimens, samples were first thinned by EDM and then thinned by grinding using pads of reduced grit size each time. Further thinning was performed by polishing with diamond suspensions of 9 μm, 3 μm and 1 μm, respectively, to produce 60-70 μm thick foils. A 3 mm diameter disk was punched out of the foil and final polishing was completed by electropolishing using a twin-jet polisher. The chemical solution used was 30% nitric acid mixed in methanol base. For insufficient thin regions for TEM observations, TEM specimens can be ion milled using the Gatan Precision Ion Polishing System (PIPS). Ion milling is typically performed at 4.5 keV and the incline angle is reduced from 4 ° to 2 ° to exploit thin areas. TEM irradiation was performed using a JEOL 2100 high resolution microscope operated at 200 kV. X-ray diffraction was performed using a PANalytical X'Pert MPD diffractometer with a Cu Ka X-ray tube and operated at 45 kV with a filament current of 40 mA. The scan was performed with 2θ of 25 ° to 95 ° combined with silicon to adjust a step size of 0.01 ° and zero angle shift of the instrument. The scans obtained were then analyzed using Rietveld analysis using Siroquant software.

After solidification, a modal structure was formed in the 50 mm thick Alloy 1 slab. The modal structure (structure # 1) is shown as a dendritic structure composed of several phases. In FIG. 6A, the backscatter SEM image shows dendritic cancer, which appears to be dark contrast, while the matrix phase is bright contrast. Note that small cast pores are shown in the SEM micrographs (black holes). TEM irradiation shows that the matrix phase is mainly austenite (γ-Fe) with lamination defects. The presence of lamination defects indicates a face centered cubic structure (austenite). TEM also suggests that other phases can be formed within the modal structure. As shown in Fig. 6C, a dark phase is found which is specified as a ferrite phase having a body centered cubic structure (α-Fe) according to the selected electron diffraction pattern. X-ray diffraction analysis showed that the modal structure of Alloy 1 contained austenite, ferrite, iron manganese compounds and some martensite (FIG. 7). In general, austenite is the dominant phase in the modal structure of alloy 1, but other factors such as cooling rate during commercial production can affect the formation of secondary phases, such as martensite of varying volume fractions.

Table 11 X-ray diffraction data of alloy 1 after solidification (modal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.583 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.876 Å Martensite Structure: Square
Space group #: 139 (I4 / mmm)
LP: a = 2.898 Å
c = 3.018 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.093 Å

Deformation of Alloy 1 with a modal structure (structure # 1, FIG. 1A) at elevated temperature leads to homogenization and refinement of the modal structure. Although hot rolling has been applied in this case, other processes, including but not limited to hot pressing, hot forging, hot extrusion, can achieve similar effects. During hot rolling, the dendrite in the modal structure is crushed and refined to cause the formation of the first homogenized modal structure (structure # 1a, FIG. 1A). Micronization during hot rolling occurs through nanophase micronization (mechanism # 1, FIG. 1A) accompanied by dynamic recrystallization. The homogenized modal structure can be gradually refined by repeatedly applying hot rolling and can lead to the formation of nanomodal structure (structure # 2, FIG. 1A). FIG. 8A shows backscatter SEM micrographs of Alloy 1 after hot rolling from 50 mm to about 1.7 mm at 1250 ° C. FIG. It can be seen that during hot rolling a block of several tens of microns is obtained from the dynamic recrystallization, and the interior of the grains is relatively smooth, indicating that the amount of defects is small. TEM also shows that sub-grains less than several hundred nanometers in size are formed, as shown in FIG. 8B. X-ray diffraction analysis shows that the nanomodal structure of Alloy 1 after hot rolling contains mainly austenite with other phases such as ferrite and iron manganese compounds as shown in FIG. 9 and Table 12.

[Table 12] X-ray diffraction data of alloy 1 after hot rolling (nano modal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.595 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.896 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.113 Å

Further deformation (i.e. cold deformation) at ambient temperature of Alloy 1 with nanomodal structure is transformed into high strength nanomodal structure (structure # 3, FIG. 1A) through dynamic nanophase reinforcement (mechanism # 2, FIG. 1A). Cause. Cold deformation can be achieved by cold rolling, tensile deformation, or other types of deformation such as punching, extrusion, stamping, and the like. During cold deformation, depending on the chemistry of the alloy, much of the austenite in the nanomodal structure is transformed into ferrite by grain refinement. 10A shows backscatter SEM micrographs of cold rolled alloy 1. FIG. Compared with the smooth grains in the nanomodal structure after hot rolling, the cold deformed grains are coarse, indicating severe plastic deformation in the grains. Depending on the chemical nature of the alloy, strain twins may be produced in some alloys, especially by cold rolling, as indicated in FIG. 10A. FIG. 10B shows a TEM micrograph of the microstructure in cold rolled alloy 1. FIG. It can be seen that in addition to the dislocations generated by the deformation, micronized grains due to phase transformation can also be found. The banded structure is associated with the deformation twin caused by cold rolling corresponding to that of FIG. 10A. X-ray diffraction shows that the high strength nanomodal structure of Alloy 1 after cold rolling contains a significant amount of ferrite phase in addition to residual austenite and iron manganese as shown in FIG. 11 and Table.

Table 13 X-ray diffraction data of alloy 1 after cold rolling (high intensity nanomodal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.588 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.871 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.102 Å

Recrystallization occurs upon heat treatment of cold strained alloy 1 having a high strength nanomodal structure (structure # 3, FIGS. 1A and 1B) that transforms into a recrystallized modal structure (structure # 4, FIG. 1B). A TEM image of Alloy 1 after annealing is shown in FIG. 12. As shown, there are equiaxed grains with sharp and linear grain boundaries in the structure, which do not have dislocations, which are characteristic of recrystallization. Depending on the annealing temperature, the size of the recrystallized grains may range from 0.5 to 50 μm. Also, as seen in electron diffraction, austenite is the dominant phase after recrystallization. Sometimes annealing twins are found within grains, but lamination defects are most often seen. Formation of lamination defects seen in TEM images is typical in the face-centered cubic crystal structure of austenite. Backscatter SEM micrographs of FIG. 13 show equiaxed recrystallized grains with a size of less than 10 μm consistent with TEM. The different contrast (dark or light contrast) of the grains seen in the SEM image suggests that the crystal orientation of the grains is irregular, in which case the contrast is mainly derived from the grain orientation. As a result, any tissue formed by the previous cold deformation is removed. X-ray diffraction shows that the recrystallized modal structure of Alloy 1 after annealing contains mainly austenite phase, including small amounts of ferrite and iron manganese compounds as shown in FIG. 14 and Table 14.

Table 14 X-ray diffraction data of Alloy 1 after annealing (recrystallized modal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.597 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.884 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.103 Å

When alloy 1 with the recrystallized modal structure (structure # 4, FIG. 1B) is deformed at ambient temperature, nanophase refinement and reinforcement (mechanism # 4, FIG. 1B) is activated to refine the high strength nanomodal structure (structure # 5). , 1b). In this case, the deformation was the result of a tensile test and the gage section of the tensile sample after the test was analyzed. FIG. 15 shows brightfield TEM micrographs of the microstructure of modified Alloy 1. FIG. Compared to matrix grains that initially have little dislocation in the recrystallized modal structure after annealing, applying a stress produces a high density of dislocations in the matrix grains. At the end of the tensile strain (tensile elongation exceeds 50%), accumulation of a large number of dislocations is found in the matrix grains. As shown in FIG. 15A, in some regions (eg, the region of the bottom of FIG. 15A), the dislocations form a cell structure, and the matrix remains austenite. In other regions, when the dislocation density is sufficiently high, transformation from austenite to ferrite (eg, upper right of 15a) is induced, resulting in significant structural refinement. FIG. 15B shows the local “pockets” of the transformed microstructure, with the selected region electron diffraction pattern corresponding to ferrite. Structural transformation into micronized high strength nanomodal structures (structure # 5, FIG. 1B) in irregularly distributed “pockets” is a feature of the steel alloys herein. 16 shows a backscatter SEM image of the refined high strength nanomodal structure. Compared to the recrystallized modal structure, the matrix grain boundaries become less pronounced and the matrix is clearly deformed. The details of the modified grain cannot be revealed by SEM, but the change caused by the deformation is enormous compared to the recrystallized modal structure demonstrated by the TEM image. X-ray diffraction shows that the refined high strength nanomodal structure of Alloy 1 after tensile deformation contains a sufficient amount of ferrite and austenite phases. Very wide peaks of the ferrite phase (α-Fe) are seen in the XRD pattern, suggesting significant refinement of the phase. Iron manganese compounds are also present. In addition, a hexagonal phase with spatial group # 186 (P6 _3mc ) was identified in the gauge portion of the tensile sample as shown in FIG. 17 and Table 15. FIG.

Table 15 X-ray diffraction data of alloys after tensile deformation (micronized high strength nanomodal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.586 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.873 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.159 Å Hexagonal Phase 1 Structure: Hexagonal
Space group #: 186 (P6 ₃ mc)
LP: a = 3.013 Å, c = 6.183 Å

This example demonstrates that the alloys listed in Table 2, including Alloy 1, represent structural evolution pathways with the novel effective mechanisms illustrated in FIGS. 1A and 1B and result in unique microstructures with nanoscale features.

Example # 2 Structural Development Pathway of Alloy 2

A 50 mm thick laboratory slab was cast from Alloy 2 and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes in the laboratory as described in the body section herein. The microstructure of the alloy at each process step was examined by SEM, TEM and X-ray analysis.

로 수행되었다. 다음에 얻어진 스캔을 Siroquant 소프트웨어를 사용하는 리트벨트 분석을 사용하여 분석하였다.For SEM irradiation, the cross section of the slab sample was polished with a small grit SiC abrasive paper and then gradually polished with a diamond media paste of up to 1 μm. The final polishing was performed using an SiO ₂ solution in 0.02 μm grit. Microstructure was examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. To prepare TEM specimens, samples were first thinned by EDM and then thinned by grinding using pads of reduced grit size each time. Further thinning was performed by polishing with diamond suspensions of 9 μm, 3 μm and 1 μm, respectively, to produce foils about 60 μm thick. A 3 mm diameter disk was punched out of the foil and final polishing was completed by electropolishing using a twin-jet polisher. The chemical solution used was 30% nitric acid mixed in methanol base. For insufficient thin regions for TEM observations, TEM specimens can be ion milled using the Gatan Precision Ion Polishing System (PIPS). Ion milling is typically performed at 4.5 keV and the incline angle is reduced from 4 ° to 2 ° to exploit thin areas. TEM irradiation was performed using a JEOL 2100 high resolution microscope operated at 200 kV. X-ray diffraction was performed using a Panalytical X'Pert MPD diffractometer with a Cu Ka X-ray tube and operated at 45 kV with a filament current of 40 mA. The scan is 25 ° to 95 ° 2 with silicon combined to adjust a step size of 0.01 ° and zero angle shift of the instrument.

Was performed. The scans obtained were then analyzed using Rietveld analysis using Siroquant software.

A modal structure (structure # 1, FIG. 1A) is formed in the slab of alloy 2 cast to a thickness of 50 mm, which is characterized by a dendritic structure. Due to the presence of the boride phase (M ₂ B), the dendritic structure is clearer than Alloy 1, in which no boride is present. 18A shows a backscatter SEM of a modal structure showing a dendritic matrix (bright contrast) with borides at the interface (dark contrast). TEM irradiation shows that the matrix phase consists of austenite (γ-Fe) with lamination defects (FIG. 18B). Similar to Alloy 1, the presence of lamination defects indicates that the matrix phase is austenite. The TEM also has a boride phase that appears dark in FIG. 18B at the interface on the austenite matrix. X-ray diffraction analysis data in FIG. 19 and Table 16 show that the modal structure contains austenite, M ₂ B, ferrite, and iron manganese compounds. Similar to alloy 1, austenite is the dominant phase of the modal structure of alloy 2, but other phases may exist depending on the chemical nature of the alloy.

TABLE 16 X-ray diffraction data of alloy 2 after solidification (modal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.577 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.850 Å M ₂ B Structure: Square
Space group #: 140 (I4 / mcm)
LP: a = 5.115 Å, c = 4.226 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.116 Å

Following the flow chart of FIG. 1A, deformation of Alloy 2 with a modal structure (structure # 1, FIG. 1A) at elevated temperature causes homogenization and refinement of the modal structure. Although hot rolling has been applied in this case, other processes, including but not limited to hot pressing, hot forging, hot extrusion, can achieve similar effects. During hot rolling, the dendrite in the modal structure is crushed and refined to cause the formation of the first homogenized modal structure (structure # 1a, FIG. 1A). Micronization during hot rolling occurs through nanophase micronization (mechanism # 1, FIG. 1A) accompanied by dynamic recrystallization. The homogenized modal structure can be gradually refined by repeatedly applying hot rolling and can lead to the formation of nanomodal structure (structure # 2, FIG. 1A). 20A shows a backscatter SEM micrograph of Hot Rolled Alloy 2. FIG. Similar to Alloy 1, the dendritic modal structure is homogenized while the boride phase is irregularly distributed in the matrix. TEM shows partial recrystallization as a result of dynamic recrystallization during hot rolling as shown in FIG. 20B. The matrix grain is on the order of 500 nm, which is finer than Alloy 1 due to the pinning effect of the boride. X-ray diffraction analysis shows that the nanomodal structure of Alloy 2 after hot rolling contains mainly austenite phase and M ₂ B along with other phases such as ferrite and iron manganese compounds as shown in FIG. 21 and Table 17.

Table 17 X-ray diffraction data of Alloy 2 after hot rolling (nanomodal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.598 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.853 Å M ₂ B Structure: Square
Space group #: 140 (I4 / mcm)
LP: a = 5.123 Å, c = 4.182 Å Iron manganese compounds Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 4.180 Å

Deformation (ie, cold deformation) of Alloy 2 with nanomodal structures but at ambient temperature results in the formation of high strength nanomodal structures (structure # 3, FIG. 1A) through dynamic nanophase reinforcement (mechanism # 2, FIG. 1A). do. Cold deformation can be achieved by cold rolling, tensile deformation, or other types of deformation such as punching, extrusion, stamping, and the like. Similarly to Alloy 2, during cold deformation, much of the austenite in the nanomodal structure is transformed into ferrite by grain refinement. 22A shows backscatter SEM micrographs of the microstructures in cold rolled alloy 2. FIG. The deformation is concentrated on the matrix around the boride phase. 22B shows a TEM micrograph of cold rolled Alloy 2. FIG. Micronized grains due to phase transformation can be found. Strain twins are less evident in the SEM image, but TEM shows that strain twins are produced after cold rolling, similar to Alloy 1. X-ray diffraction showed that the high-strength nanomodal structure of Alloy 2 after cold rolling, in addition to a new hexagonal phase with M ₂ B, residual austenite and space group # 186 (P6 _3mc ), as shown in FIG. 23 and Table 18 It shows that it contains a significant amount of ferrite phase.

Table 18 X-ray diffraction data of alloy 2 after cold rolling (high intensity nanomodal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.551 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.874 Å M ₂ B Structure: Square
Space group #: 140 (I4 / mcm)
LP: a = 5.125 Å, c = 4.203 Å Hexagonal phase Structure: Hexagonal
Space group #: 186 (P6 ₃ mc)
LP: a = 2.962 Å, c = 6.272 Å

Recrystallization occurs upon annealing of cold strained alloy 2 with a high strength nanomodal structure (structure # 3, Figures 1A and 1B) that transforms into a recrystallized modal structure (structure # 4, Figure 1B). The recrystallized microstructure of Alloy 2 after annealing is shown in the TEM image of FIG. 24. As shown, there are equiaxed grains with sharp and linear grain boundaries in the structure, which do not have dislocations, which are characteristic of recrystallization. The size of the recrystallized grains is generally less than 5 μm due to the pinning effect on the boride, but larger grains are possible at higher annealing temperatures. Electron diffraction also shows that austenite is the dominant phase after recrystallization and that there is a lamination bond in austenite, as shown in FIG. 24B. The formation of lamination defects also indicates the formation of faceted cubic austenite phases. Backscatter SEM micrographs of FIG. 25 show equiaxed recrystallized grains of size less than 5 μm with irregularly distributed boride phases. The different contrast (dark or light contrast) of the grains seen in the SEM image suggests that the crystal orientation of the grains is irregular, in which case the contrast is mainly derived from the grain orientation. As a result, any tissue formed by the previous cold deformation is removed. X-ray diffraction shows that the recrystallized modal structure of Alloy 2 after annealing is mainly austenite with M ₂ B, a small amount of ferrite, and a hexagonal phase with space group # 186 (P6 _3mc ) as shown in FIG. 26 and Table 19 It shows that it contains a phase.

TABLE 19 X-ray diffraction data of Alloy 2 after annealing (recrystallized modal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.597 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.878 Å M ₂ B Structure: Square
Space group #: 140 (I4 / mcm)
LP: a = 5.153 Å, c = 4.170 Å Hexagonal phase Structure: Hexagonal
Space group #: 186 (P6 _3mc )
LP: a = 2.965 Å, c = 6.270 Å

Deformation of the recrystallized modal structure (structure # 4, FIG. 1B) leads to the formation of a refined high strength nanomodal structure (structure # 5, FIG. 1B) through nanophase refinement and reinforcement (mechanism # 4, FIG. 1B). In this case, the deformation was the result of a tensile test and the gauge portion of the tensile sample after the test was analyzed. FIG. 27 shows a micrograph of the microstructure of modified Alloy 2. FIG. Similar to Alloy 1, the initial dislocation-free matrix grains of the recrystallized modal structure after annealing are filled with high-density dislocations upon application of stress, and the accumulation of dislocations in some grains activates phase transformation of ferrite from austenite, Cause substantial refinement. As shown in FIG. 27A, micronized grains of 100-300 nm size are in local “pockets” where transformation from austenite to ferrite has occurred. Structural transformation of the high-strength nanomodal structure (structure # 5, FIG. 1B) micronized within the “pocket” of the matrix grains is a feature of the steel alloys herein. 27B shows a backscatter SEM image of the refined high strength nanomodal structure. Similarly, the grain boundaries of the matrix grains become less apparent after the matrix is deformed. X-ray diffraction shows that the four phases remain as in the recrystallized modal structure but a significant amount of austenite has been transformed into ferrite. The transformation resulted in the formation of a refined high strength nanomodal structure of alloy 2 after tensile deformation. Very wide peaks of the ferrite phase (α-Fe) are seen in the XRD pattern, suggesting significant refinement of the phase. As in Alloy 1, a new hexagonal phase with space group # 186 (P6 _3mc ) was identified in the gauge portion of the tensile sample as shown in FIG. 28 and Table 20.

TABLE 20 X-ray diffraction data of Alloy 2 after tensile deformation (micronized high strength nanomodal structure)

Verified award On Pinterest γ-Fe Structure: Cubic
Space group #: 225 (Fm3m)
LP: a = 3.597 Å α-Fe Structure: Cubic
Space group #: 229 (Im3m)
LP: a = 2.898 Å M ₂ B Structure: Square
Space group #: 140 (I4 / mcm)
LP: a = 5.149 Å, c = 4.181 Å Hexagonal phase Structure: Hexagonal
Space group #: 186 (P6 _3mc )
LP: a = 2.961 Å, c = 6.271 Å

This example demonstrates that the alloys listed in Table 2, including Alloy 2, represent structural evolution pathways with the mechanisms illustrated in FIGS. 1A and 1B and result in unique microstructures with nanoscale features.

Example # 3 Tensile Properties at Each Process Step

Slabs with a thickness of 50 mm were laboratory cast with the alloys listed in Table 21 according to the atomic ratios provided in Table 2, and 10 minutes at 850 ° C. as described in the hot rolling, cold rolling and body parts herein of the laboratory treatment. Annealed for a while. Tensile properties were measured on Instron 3369 mechanical test frames using Instron's Bluehill control software at each process step. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer.

The alloys were weighed into charges in the range of 3,000 to 3,400 grams using commercially available ferroditive powders with known chemical properties and impurity contents according to the atomic ratios of Table 2. The charge was charged into a zirconia coated silica crucible installed in an Indutherm VTC800V vacuum tilt casting machine. The casting machine was then evacuated to cast and melt chambers and backfilled with argon several times before casting to prevent oxidation of the melt. The melt was heated with a 14 kHz RF induction coil until it melted completely for about 5.25 to 6.5 minutes, depending on the alloy composition and the mass of the charge. After melting of the last solid was observed, it was further heated for 30 to 45 seconds to provide overheating to ensure homogenization of the melt. The casting machine then emptied the melting and casting chamber and tilted the crucible to pour the melt into a 50 mm thick, 75 to 80 mm wide, and 125 mm deep channel of the water cooled copper die. The melt was cooled under vacuum for 200 seconds and then the chamber was filled with argon to atmospheric pressure. Tensile specimens were cut from the slab as it was cast by wire EDM and a tensile test was performed. The results of the tensile test are shown in Table 21. As can be seen, the maximum tensile strength of the alloy of the present specification in the cast state ranges from 411 to 907 MPa. Tensile elongation is in the range of 3.7 to 24.4%. Yield strength is measured in the range of 144 to 514 MPa.

Prior to hot rolling, the lab cast slab was charged by heating into a Lucifer EHS3GT-B18 furnace. The set point of the furnace ranges from 1000 ° C. to 1250 ° C. depending on the alloy melting point. This was soaked for 40 minutes before hot rolling to ensure that the slab reached the target temperature. Between the hot rolling and the hot rolling the slab was returned to the furnace and reheated for 4 minutes. The preheated slab was pushed from the tunnel furnace into the Fenn Model 061 2 high pressure smoke. The 50 mm casting was hot rolled through a rolling mill over five to eight passes followed by air cooling as defined in the first campaign of hot rolling. The slab thickness was reduced from 80.4 to 87.4% after this campaign. After cooling, the obtained sheet sample was cut to 190 mm in length. These cuts were hot rolled through a rolling mill over a further three passes with a reduction of 73.1 to 79.9% to obtain a final thickness of 2.1 to 1.6 mm. Details regarding the hot rolling conditions for each alloy herein are provided in Table 22. Tensile specimens were cut from the hot rolled sheet by wire EDM and a tensile test was performed. The results of the tensile test are shown in Table 22. After hot rolling, the maximum tensile strength of the alloys herein is in the range of 921-1413 MPa. Tensile elongation is in the range of 12.0 to 77.7%. Yield strength is measured in the range of 264 to 574 MPa. See structure 2 of FIG. 1A.

The sheet obtained after hot rolling was media blasted with aluminum oxide to remove the rolling scale, which was then cold rolled on a Fenn Model 061 2 high pressure steamer. Cold rolling takes a number of passes to reduce the thickness of the sheet to a target thickness of generally 1.2 mm. The hot rolled sheet was fed into the rolling mill continuously decreasing the roll gap until the minimum gap was reached. If the material did not reach the specification target, an additional pass was used with minimum clearance until the target thickness was reached. Cold rolling conditions with pass recovery for each alloy in this specification are reported in Table 23. Tensile specimens were cut from the cold rolled sheet by wire EDM and tensile tests were performed. The results of the tensile test are shown in Table 23. Cold rolling results in sufficient reinforcement with maximum tensile strength in the range of 1356 to 1831 MPa. The tensile elongation of the alloy of the present specification in the cold rolled state is in the range of 1.6 to 32.1%. Yield strength is measured in the range of 793-1645 MPa. It is expected that greater maximum tensile and yield strengths can be achieved in the alloys herein by greater cold rolling reduction (greater than 40%) (in our case limited due to the ability of the mill's mill). When using a larger rolling force, it is expected that the maximum tensile strength can be increased by at least 2000 MPa and the yield strength by at least 1800 MPa.

Tensile specimens were cut from cold rolled sheet samples by wire EDM and annealed at 850 ° C. for 10 minutes in a Lucifer 7HT-K12 box furnace. Samples were removed from the furnace at the end of the cycle and cooled to room temperature in air. The results of the tensile test are shown in Table 24. As can be seen, recrystallization during annealing of the alloys of the present disclosure yields a combination of properties of maximum tensile strength in the range of 939 to 1424 MPa and tensile elongation of 15.8 to 77.0%. Yield strength is measured in the range of 420 to 574 MPa. 29-31 show the data plotted at each process step for Alloy 1, Alloy 13, and Alloy 17, respectively.

[Table 21] Tensile Properties of Alloys in Cast State

alloy Yield strength (MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 289 527 10.4 288 548 9.3 260 494 8.4 Alloy 2 244 539 10.4 251 592 11.6 249 602 13.1 Alloy 13 144 459 4.6 156 411 4.5 163 471 5.7 Alloy 17 223 562 24.4 234 554 20.7 235 585 23.3 Alloy 24 396 765 8.3 362 662 5.7 404 704 7.0 Alloy 25 282 668 5.1 329 753 5.0 288 731 5.5 Alloy 25 471 788 4.1 514 907 6.0 483 815 3.7 Alloy 27 277 771 3.7 278 900 4.9 267 798 4.5 Alloy 34 152 572 11.1 168 519 11.6 187 545 12.9 Alloy 35 164 566 15.9 172 618 16.6 162 569 16.4

Table 22 Tensile Properties of Alloys in Hot Rolled State

alloy Condition 1st campaign reduction 2nd campaign reduction Yield strength
(MPa) Tensile strength (MPa) Tensile elongation
(%) Alloy 1 Hot Rolled95.2% 80.5%,
6 pass 75.1%,
3 pass 273 1217 50.0 264 1216 52.1 285 1238 52.7 Alloy 2 Hot Rolled96.6% 87.4%,
7 pass 73.1%,
3 pass 480 1236 45.3 454 1277 41.9 459 1219 48.2 Alloy 13 Hot Rolled96.0% 81.1%,
6 pass 79.8%,
3 pass 287 1116 18.8 274 921 15.3 293 1081 19.3 Alloy 17 Hot Rolled96.1% 81.2%,
6 pass 79.1%,
3 pass 392 947 73.3 363 949 74.8 383 944 77.7 Alloy 24 Hot rolled, 96.2% 81.1%,
6 pass 79.9%
3 pass 519 1176 21.4 521 1088 18.2 508 1086 17.9 Alloy 25 Hot Rolled96.1% 81.0%,
6 pass 79.4%,
3 pass 502 1105 12.4 524 1100 12.3 574 1077 12.0 Alloy 27 Hot rolled, 95.9% 80.4%,
6 pass 78.9%,
3 pass 508 1401 20.9 534 1405 22.4 529 1413 19.7 Alloy 34 Hot rolled, 96.2% 80.7%,
6 pass 80.1%,
3 pass 346 1188 56.5 323 1248 58.7 303 1230 53.4 Alloy 35 Hot rolled, 96.1% 80.8%,
6 pass 79.9%,
3 pass 327 1178 63.3 317 1170 61.2 305 1215 59.6

Table 23 Tensile Properties of Alloys in Cold Rolled State

alloy Condition Yield strength (MPa) Tensile strength (MPa) Tensile Elongation (%) Alloy 1 Cold rolled 20.3%,
4 pass 798 1492 28.5 793 1482 32.1 Cold rolled 39.6%,
29 passes 1109 1712 21.4 1142 1726 23.0 1203 1729 21.2 Alloy 2 Cold rolled 28.5%,
5 pass 966 1613 13.4 998 1615 15.4 1053 1611 20.6 Cold rolled39.1%,
19 passes 1122 1735 20.3 1270 1744 18.3 Alloy 13 Cold rolled 36.0%,
24 passes 1511 1824 9.5 1424 1803 7.7 1361 1763 5.1 Alloy 17 Cold rolled 38.5%,
8 pass 1020 1357 24.2 1007 1356 24.9 1071 1357 24.9 Alloy 24 Cold rolled 38.2%,
23 passes 1363 1584 1.9 1295 1601 2.5 1299 1599 3.0 Alloy 25 Cold rolled 38.0%,
42 passes 1619 1761 1.9 1634 1741 1.7 1540 1749 1.6 Alloy 27 Cold rolled 39.4%,
40 passes 1632 1802 2.7 1431 1804 4.1 1645 1831 4.1 Alloy 34 Cold rolled 35.%,
14 passes 1099 1640 14.7 840 1636 17.5 1021 1661 18.5 Alloy 35 Cold rolled 35.5%,
12 passes 996 1617 23.8 1012 1614 24.5 1020 1616 23.3

Table 24 Tensile Properties of Alloys in the Annealed State

alloy Yield strength (MPa) Tensile strength (MPa) Tensile Elongation (%) Alloy 1 436 1221 54.9 443 1217 56.0 431 1216 59.7 Alloy 2 438 1232 49.7 431 1228 49.8 431 1231 49.4 484 1278 48.3 485 1264 45.5 479 1261 48.7 Alloy 13 441 1424 41.7 440 1412 41.4 429 1417 42.7 Alloy 17 420 946 74.6 421 939 77.0 425 961 74.9 Alloy 24 554 1151 23.5 538 1142 24.3 562 1151 24.3 Alloy 25 500 1274 16.0 502 1271 15.8 483 1280 16.3 Alloy 27 538 1385 20.6 574 1397 20.9 544 1388 21.8 Alloy 27 467 1227 56.7 476 1232 52.7 462 1217 51.6 Alloy 27 439 1166 56.3 438 1166 59.0 440 1177 58.3

This example shows that the unique mechanisms and structural pathways shown in FIGS. 1A and 1B can lead to the development of third generation AHSS by varying the structure and consequent properties of the alloys herein.

Example # 4 Cyclic Reversibility in Cold Rolling and Recrystallization

A slab with a thickness of 50 mm was laboratory cast in alloy 1 and alloy 2 according to the atomic ratios given in Table 2, with a final thickness of 2.31 mm for alloy 1 sheet and 2.35 mm for alloy 2 sheet. The sheet was hot rolled into a sheet having a final thickness of. Casting and hot rolling procedures are described in the text of this application. The hot rolled sheet obtained from each alloy was used to explain the circulation structure / characteristic reversibility through the cold rolling / annealing cycle.

The 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 in each cycle are reported in Table 25. Annealing was carried out at 850 ° C. for 10 minutes after each cold rolling. Tensile specimens were cut from the sheet in the initial hot rolled state and at each cycling step. Tensile properties were measured on an Instron 3369 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer.

The tensile test results for Alloy 1 and Alloy 2 are plotted in FIG. 32, which shows that by cold rolling a significant strengthening was obtained for both alloys in each cycle with an average maximum tensile strength of 1500 MPa in Alloy 1 and 1580 MPa in Alloy 2. Shows. Both cold rolled alloys show a loss of malleability compared to the hot rolled state. However, the annealing after cold rolling in each cycle causes the tensile property recovery to the same level as the high malleability.

Tensile properties for each tested sample are reported in Table 26 and Table 27 for each of Alloy 1 and Alloy 2. As can be seen, Alloy 1 has a maximum tensile strength of 1216 to 1238 MPa, a malleability of 50.0 to 52.7% and a yield strength of 264 to 285 MPa in the hot rolled state. In the cold rolled state, the maximum tensile strength was measured in the range of 1482 to 1517 MPa in each cycle. Malleability was consistently observed in the 28.5 to 32.8% range with significantly higher yield strengths of 718 to 830 MPa compared to that of the hot rolled state. Annealing in each cycle returned to malleability ranging from 47.7 to 59.7% with a maximum tensile strength of 1216 to 1270 MPa. The yield strength after cold rolling and annealing was measured in the range of 431 to 515 MPa lower than that after cold rolling but higher than that of the initial hot rolled state.

Similar results with characteristic reversibility between cold rolled and annealed materials through cycling were observed for Alloy 2 (FIG. 32B). In the initial hot rolled state, Alloy 2 has a maximum tensile strength of 1219 to 1277 MPa with a malleability of 41.9 to 48.2% and a yield strength of 454 to 480 MPa. Cold rolling in each cycle results in material reinforcement up to a maximum tensile strength of 1553 to 1598 MPa and a decrease in malleability in the range of 20.3 to 24.1%. Yield strength was measured at 912-1126 MPa. After annealing in each cycle, Alloy 2 has a maximum tensile strength of 1231 to 1281 MPa and a malleability of 46.9 to 53.5%. The yield strength of Alloy 2 after cold rolling and annealing in each cycle is similar to that in the hot rolled state and ranges from 454 to 521 MPa.

Table 25 Sample Thickness and Cycle Reduction in Cold Rolling Step

alloy Rolling cycle Initial thickness (mm) Final thickness (mm) % Cycle reduction Alloy 1 One 2.35 1.74 26.0 2 1.74 1.32 24.1 3 1.32 1.02 22.7 Alloy 2 One 2.31 1.85 19.9 2 1.85 1.51 18.4 3 1.51 1.22 19.2

Table 26 Tensile Properties of Alloy 1 Through Cold Rolling / Annealing Cycles

characteristic Hot rolled 1st cycle 2nd cycle 3rd cycle Cold rolled Annealed Cold rolled Annealed Cold rolled Annealed Tensile strength (MPa) 1217 1492 1221 1497 1239 1517 1270 1216 1482 1217 1507 1269 1507 1262 1238 * 1216 1503 1260 1507 1253 Yield strength (MPa) 273 798 436 775 487 820 508 264 793 443 718 466 796 501 285 * 431 830 488 809 515 Tensile Elongation (%) 50.0 28.5 54.9 32.8 57.5 32.1 50.5 52.1 32.1 56.0 29.4 52.5 30.2 47.7 52.7 * 59.7 30.9 55.8 30.5 55.5

* Specimen deviates from grip / data not available

Table 27 Tensile Properties of Alloy 2 through Cold Rolling / Annealing Cycles

characteristic Hot rolled 1st cycle 2nd cycle 3rd cycle Cold rolled Annealed Cold rolled Annealed Cold rolled Annealed Tensile strength (MPa) 1236 1579 1250 1553 1243 1596 1231 1277 * 1270 1568 1255 1589 1281 1219 * 1240 1566 1242 1598 1269 Yield strength (MPa) 480 1126 466 983 481 1006 475 454 * 468 969 521 978 507 459 * 454 912 497 1011 518 Tensile Elongation (%) 45.3 20.3 53.0 24.1 51.1 22.3 46.9 41.9 * 51.2 23.1 52.3 23.2 53.5 48.2 * 51.1 21.6 49.9 21.0 47.9

* Specimen Departs from Grip / No Data Available This example shows a modal in which a high strength nanomodal structure (structure # 3, FIG. 1A) formed from the alloys listed in Table 2 after cold rolling is recrystallized by annealing and recrystallized. It is shown that a structure (structure # 4, FIG. 1B) can be created. This structure may be further deformed through cold rolling or other cold deformation methods to lead to the formation of micronized high strength nanomodal structures (structure # 5, FIG. 1B) through nanophase refinement and reinforcement (mechanism # 4, FIG. 1B). . The refined high-strength nanomodal structure (structure # 5, FIG. 1B) can then be recrystallized and the process can begin again with complete structure / characteristic reversibility over multiple cycles. The reversible capability of the mechanism allows the production of finer specifications that are important for weight reduction and recovery of properties after damage due to deformation when using AHSS.

Example # 5 flexing ability

Slabs with a thickness of 50 mm were laboratory cast with the selected alloys listed in Table 28 according to the atomic ratios given in Table 2, and hot-rolled, cold rolled by laboratory treatment and 10 at 850 ° C. as described in the text section herein. Annealed for minutes. The flexural response of the alloys of this disclosure was evaluated using sheets obtained from each alloy with a final thickness of about 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B).

Flexural testing was performed using an Instron 5984 tensile test platform with Instron W-6810 guided flexural test fixture in accordance with the specifications outlined in ISO 7438 International Standard Metal Material-Flexation Test (International Organization for Standardization, 2005). The specimens were cut to dimensions of 20 mm x 55 mm x sheet thickness by wire EDM. No special edge preparation for the sample was performed. Flexural testing was performed using an Instron 5984 tensile test platform with Instron W-6810 guided flexural test fixture. Flexural testing was performed according to the specifications outlined in ISO 7438 International Organization for Standardization (2005).

The test was performed by placing the specimen on the fixture support as shown in FIG. 33 and pushing it with the former.

The distance between supports (l) was fixed as follows according to ISO 7438 during the test:

Prior to bending, both sides of the specimen were lubricated with "3 in 1" oil to reduce friction with the test fixture. This test was performed using a 1 mm diameter former. The former was pressed at various angles up to 180 ° or downwards in the middle of the support until cracks appeared. Flexural forces were applied slowly to allow free plastic flow of the material. Displacement rate was calculated based on the span gap of each test in order to make the angular rate constant and applied accordingly.

The absence of visible cracks without the use of magnifying means was considered evidence that the specimens survived the flexural test. When a crack was detected, the bending angle was measured manually using a digital protractor at the bottom of the bending portion. The specimen was then removed from the fixture and inspected for cracks outside the radius of the bend. The initiation of the crack could not be determined decisively from the force-displacement curve, but instead was easily determined by direct observation with illumination from the flashlight.

The results of the bending reaction of the alloys herein are reported in Table 28, including the initial sheet thickness, the ratio of the former radius to the sheet thickness (r / t) and the maximum bending angle before cracking. All alloys listed in Table 28 showed no cracks at 90 ° bend angles. Most of the alloys herein have the ability to bend at 180 ° without cracking. An example of a sample from alloy 1 after a 180 ° bend test is shown in FIG. 34.

TABLE 7 Table 7 Bending Test Results for Selected Alloys

alloy Former diameter
(mm) Thickness (mm) r / t Bending angle (°) Alloy 1 0.95 1.185 0.401 180 1.200 0.396 180 1.213 0.392 180 1.223 0.388 180 1.181 0.402 180 1.187 0.400 180 1.189 0.399 180 1.206 0.394 180 Alloy 2 0.95 1.225 0.388 180 1.230 0.386 180 1.215 0.391 180 1.215 0.391 180 1.215 0.391 180 1.224 0.388 180 1.208 0.393 180 1.208 0.393 180 Alloy 3 0.95 1.212 0.392 180 1.186 0.401 180 1.201 0.396 180 Alloy 4 0.95 1.227 0.387 180 1.185 0.401 180 1.187 0.400 180 Alloy 5 0.95 1.199 0.396 110 1.196 0.397 90 Alloy 6 0.95 1.259 0.377 160 1.202 0.395 165 1.206 0.394 142 Alloy 7 0.95 1.237 0.384 104 1.236 0.384 90 Alloy 9 0.95 1.278 0.372 180 1.197 0.397 180 1.191 0.399 180 Alloy 10 0.95 1.226 0.387 180 1.208 0.393 100 1.208 0.393 180 1.205 0.394 180 Alloy 11 0.95 1.240 0.383 180 1.214 0.391 180 1.205 0.394 180 Alloy 12 0.95 1.244 0.382 180 1.215 0.391 180 1.205 0.394 180 Alloy 13 0.95 1.222 0.389 180 1.191 0.399 180 1.188 0.400 180 Alloy 14 0.95 1.239 0.383 180 1.220 0.389 180 1.214 0.391 180 Alloy 15 0.95 1.247 0.381 180 1.224 0.388 180 1.224 0.388 180 Alloy 16 0.95 1.244 0.382 180 1.224 0.388 180 1.199 0.396 180 Alloy 17 0.95 1.233 0.385 180 1.213 0.392 180 1.203 0.395 180 Alloy 18 0.95 1.222 0.389 160 1.218 0.390 135 Alloy 19 0.95 1.266 0.375 180 1.243 0.382 180 1.242 0.382 180 Alloy 20 0.95 1.242 0.382 180 1.222 0.389 180 1.220 0.389 180 Alloy 21 0.95 1.255 0.378 180 1.228 0.387 180 1.229 0.386 180 Alloy 22 0.95 1.240 0.383 180 1.190 0.399 180 1.190 0.399 180 Alloy 23 0.95 1.190 0.399 180 1.199 0.396 180 1.193 0.398 180 Alloy 28 0.95 1.222 0.389 180 1.206 0.394 180 1.204 0.395 180 Alloy 29 0.95 1.219 0.390 180 1.217 0.390 180 1.206 0.394 180 Alloy 30 0.95 1.215 0.391 180 1.212 0.392 175 1.200 0.396 180 Alloy 31 0.95 1.211 0.392 150 1.209 0.393 131 Alloy 32 0.95 1.222 0.389 180 1.221 0.389 180 1.210 0.393 180

In order to fabricate complex parts for automotive and other applications, AHSS must exhibit both bulk sheet formability and edge sheet formability. This example demonstrates the good bulk sheet formability of the alloys of Table 2 through flexural testing.

Example # 6 Punched Edge to EDM Cut Tensile Properties

Slabs with a thickness of 50 mm were laboratory cast with selected alloys listed in Table 2 according to the atomic ratios given in Table 2, and hot rolled, cold rolled and annealed for 10 minutes at 850 ° C. as described herein by laboratory treatment. It was. The tensile specimens were cut by wire discharge machining (wire-EDM) using sheets obtained from each alloy with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) (which does not compromise mechanical properties). The impact of edge damage on alloy properties was assessed by indicating a control situation or relative lack of shear and edge formation) and by punching (to determine the loss of mechanical properties due to shear). It should be understood that shearing (application of stress on the same plane as the material cross section) may be generated by a number of processing options such as piercing, drilling, cutting or cropping (cutting of the ends of certain metal parts).

Tensile tests of ASTM E8 shapes were prepared using both wire EDM cutting and punching. Tensile properties were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Tensile data for selected alloys are shown in Table 29 and shown in FIG. 35A. A reduction in properties is observed for all alloys tested, but the level of reduction varies considerably depending on the chemical nature of the alloy. Table 30 summarizes the comparison of the malleability of the malleable and punched samples to the wire EDM cut samples. In FIG. 35B the corresponding tensile curve for the selected alloy is shown and shows mechanical behavior as a function of austenite stability. For the alloys selected herein, the austenite stability is highest in alloy 12 showing high malleability and lowest in alloy 13 showing high strength. Correspondingly, alloy 12 demonstrated the lowest loss of malleability (29.7% vs. 60.5%, Table 30) in punched specimen to EDM cuts, while alloy 13 showed the highest loss of malleability in punched specimen to EDM cuts (5.2%). Versus 39.1%, Table 30). High edge damage occurs in specimens punched from alloys with low austenite stability.

TABLE 8 Tensile Properties of Punched vs. EDM Cut Specimens from Selected Alloys

alloy Cutting way Yield strength
(MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 EDM cutting 392 1310 46.7 397 1318 45.1 400 1304 49.7 Punched 431 699 9.3 430 680 8.1 422 656 6.9 Alloy 2 EDM cutting 434 1213 46.4 452 1207 46.8 444 1199 49.1 Punched 491 823 14.4 518 792 11.3 508 796 11.9 Alloy 9 EDM cutting 468 1166 56.1 480 1177 52.4 475 1169 56.9 Punched 508 1018 29.2 507 1007 28.6 490 945 23.3 Alloy 11 EDM cutting 474 1115 64.4 464 1165 62.5 495 1127 62.7 Punched 503 924 24.6 508 964 28.0 490 921 25.7 Alloy 12 EDM cutting 481 1094 54.4 479 1128 64.7 495 1126 62.4 Punched 521 954 27.1 468 978 30.7 506 975 31.2 Alloy 13 EDM cutting 454 1444 39.5 450 1455 38.7 Punched 486 620 5.0 469 599 6.3 483 616 4.5 Alloy 14 EDM cutting 484 1170 58.7 489 1182 61.2 468 1188 59.0 Punched 536 846 17.0 480 816 18.4 563 870 17.5 Alloy 18 EDM cutting 445 1505 37.8 422 1494 37.5 Punched 478 579 2.4 469 561 2.6 463 582 2.9 Alloy 21 EDM cutting 464 1210 57.6 499 1244 49.0 516 1220 54.5 Punched 527 801 11.3 511 806 12.6 545 860 15.2 Alloy 24 EDM cutting 440 1166 31.0 443 1167 32.0 455 1176 31.0 Punched 496 696 5.0 463 688 5.0 440 684 4.0 Alloy 25 EDM cutting 474 1183 15.8 470 1204 17.0 485 1223 17.4 Punched 503 589 2.1 517 579 0.8 497 583 2.1 Alloy 26 EDM cutting 735 1133 20.8 742 1109 19.0 Punched 722 898 3.4 747 894 2.9 764 894 3.1 Alloy 27 EDM cutting 537 1329 19.3 513 1323 21.4 480 1341 20.8 Punched 563 624 4.3 568 614 3.3 539 637 4.3 Alloy 34 EDM cutting 460 1209 54.7 441 1199 54.1 475 1216 52.9 Punched 489 828 15.4 486 811 14.6 499 813 14.8 Alloy 35 EDM cutting 431 1196 50.6 437 1186 52.0 420 1172 54.7 Punched 471 826 19.9 452 828 19.7 482 854 19.7

Table 30 Tensile Elongation of Specimen for Different Cutting Methods

alloy Average Tensile Elongation (%) Loss of tensile elongation
(E2 / E1) EDM cutting (E1) Punched (E2) at least maximum Alloy 1 47.2 8.1 0.14 0.21 Alloy 2 47.4 12.5 0.23 0.31 Alloy 9 55.1 27.0 0.41 0.56 Alloy 11 63.2 26.1 0.38 0.45 Alloy 12 60.5 29.7 0.42 0.57 Alloy 13 39.1 5.2 0.11 0.16 Alloy 14 59.7 17.7 0.28 0.31 Alloy 18 37.6 2.6 0.06 0.08 Alloy 21 53.7 13.0 0.20 0.31 Alloy 24 31.3 4.7 0.13 0.16 Alloy 25 16.7 1.7 0.05 0.13 Alloy 26 31.3 4.7 0.14 0.18 Alloy 27 20.5 4.0 0.15 0.22 Alloy 34 53.9 14.9 0.27 0.29 Alloy 35 52.4 19.8 0.36 0.39

As can be seen in Table 30, the EDM cut was considered representative of the optimum mechanical properties of the particular alloy without sheared edges and was processed up to the point taking structure # 4 (recrystallized modal structure). Thus, a sample with sheared edges due to punching exhibits a significant degradation of the malleability reflected by the tensile elongation measurement of the punched sample having an ASTM E8 shape. In the case of Alloy 1, the tensile elongation is initially 47.2%, and then decreases to 8.2% (a drop of 82.8%). The decrease in malleability from punching to EDM cleavage (E2 / E1) ranges from 0.57 to 0.05.

Edge states after punching and EDM cutting were analyzed by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc. A typical appearance of the specimen edge after EDM cutting for Alloy 1 is shown in FIG. 36A. The EDM cutting method minimizes damage to the cut edges and enables the measurement of the tensile properties of the material without the harmful edge effect. In wire-EDM cutting, material is removed from the edges by a series of rapidly repeated current discharges / sparks, and the paths form the edges without significant deformation or edge damage. The appearance of the sheared edge after punching is shown in FIG. 36B. Significant damage of the edges occurs in the fracture area that undergoes severe deformation during punching, leading to structural transformation of the shear affected area to the refined high strength nanomodal structure with limited malleability (FIG. 37B) while recrystallized modal near the EDM cut edge The structure was observed (FIG. 37A).

This example shows that in the case of wire-EDM cutting, the tensile properties are measured at a relatively higher level compared to after the punching. In contrast to EDM cutting, punching in the tensile test results in significant edge damage resulting in lower tensile properties. Relatively excessive plastic deformation of the sheets of the alloys of this disclosure during punching leads to structural transformations into reduced high malleability, high-strength nanomodal structures (structure # 5, FIG. 1B), premature cracking at edges and relatively low For example, a decrease in elongation and tensile strength). The magnitude of this reduction in tensile properties has been observed to depend on the alloy's chemistry with respect to austenite stability.

Example # 7 Tensile Properties and Recovery of Punched Edge to EDM Cut

Slabs of 50 mm thickness according to the atomic ratios provided in Table 2 were laboratory cast from the selected alloys reported in Table 31 and annealed for 10 minutes at hot rolling, cold rolling and 850 ° C. as described herein by laboratory treatment. . Sheets obtained from each alloy with a final thickness of 1.2 mm and recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate edge damage recovery by annealing punched tensile specimens. In the broad context of the present invention, annealing can be accomplished in a variety of ways, including but not limited to furnace heat treatment, induction heat treatment and / or laser heat treatment.

Tensile specimens of ASTM E8 shape were prepared using both wire EDM cutting and punching. The parts of the punched tensile specimens were recovered annealed at 850 ° C. for 10 minutes and then air cooled to confirm the ability to recover the properties lost by punching and shear damage. Tensile properties were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Tensile test results for the selected alloys are provided in Table 31, and are shown in FIG. 38, showing the recovery of significant mechanical properties of the punched samples after annealing.

For example, for the indicated alloy 1 EDM cut with a tensile test sample, the tensile elongation average value is about 47.2%. As noted above, in the case of punched and thus sheared edges, tensile testing of samples with such edges results in a significant reduction in this elongation value, ie mechanism # 4 and micronized high strength nanomodal structures (structure # 5, The average value of only 8.1% was due to the formation of FIG. 1 b), which is mainly present at the edges where shearing occurred but is reflected in the bulk properties measurement in the tensile test. However, upon annealing representative of the conversion to mechanism # 3 and structure # 4 (recrystallized modal structure, FIG. 1B) of FIG. 1B, tensile elongation characteristics are restored. In the case of Alloy 1, the tensile elongation reverts to an average value of about 46.2%. An exemplary tensile stress-strain curve for annealed and unannealed punched specimens is shown in FIG. 39. Table 32 provides a summary of average tensile properties, average losses, and average gain tensile elongation. Individual losses and gains are spreads greater than the average loss. Thus, in the context of this disclosure, the alloy of the present disclosure having an initial value (E ₁₎ of the tensile elongation when the front end may represent a reduction in the elongation property to the value of E _2, wherein E ₂ = (0.0.57 to 0.05) (E ₁ ). Next, upon application of Mechanism # 3, which is preferably achieved by heating / annealing in the temperature range of 450 ° C. to T _m , depending on the chemical properties of the alloy, the value of E ₂ is the elongation value E ₃ = (0.48 to 1.21 Is recovered to (E ₁ ).

TABLE 10 Tensile Properties of Punched and Annealed Specimens from Selected Alloys

alloy Cutting way Yield strength
(MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 EDM cutting 392 1310 46.7 397 1318 45.1 400 1304 49.7 Punched 431 699 9.3 430 680 8.1 422 656 6.9 Punched and annealed 364 1305 43.6 364 1315 47.6 370 1305 47.3 Alloy 2 EDM cutting 434 1213 46.4 452 1207 46.8 444 1199 49.1 Punched 491 823 14.4 518 792 11.3 508 796 11.9 Punched and annealed 432 1205 50.4 426 1191 50.7 438 1188 49.3 Alloy 9 EDM cutting 468 1166 56.1 480 1177 52.4 475 1169 56.9 Punched 508 1018 29.2 507 1007 28.6 490 945 23.3 Punched and annealed 411 1166 59.0 409 1174 52.7 418 1181 55.6 Alloy 11 EDM cutting 474 1115 64.4 464 1165 62.5 495 1127 62.7 Punched 503 924 24.6 508 964 28.0 490 921 25.7 Punched and annealed 425 1128 64.5 429 1117 57.1 423 1140 54.3 Alloy 12 EDM cutting 481 1094 54.4 479 1128 64.7 495 1126 62.4 Punched 521 954 27.1 468 978 30.7 506 975 31.2 Punched and annealed 419 1086 65.7 423 1085 63.0 415 1100 53.8 Alloy 13 EDM cutting 454 1444 39.5 450 1455 38.7 Punched 486 620 5.0 469 599 6.3 483 616 4.5 Punched and annealed 397 1432 41.4 397 1437 37.4 404 1439 40.3 Alloy 14 EDM cutting 484 1170 58.7 489 1182 61.2 468 1188 59.0 Punched 536 846 17.0 480 816 18.4 563 870 17.5 Punched and annealed 423 1163 58.3 412 1168 55.9 415 1177 51.5 Alloy 18 EDM cutting 445 1505 37.8 422 1494 37.5 Punched 478 579 2.4 469 561 2.6 463 582 2.9 Punched and annealed 398 1506 36.3 400 1502 40.3 392 1518 35.4 Alloy 21 EDM cutting 464 1210 57.6 499 1244 49.0 516 1220 54.5 Punched 527 801 11.3 511 806 12.6 545 860 15.2 Punched and annealed 409 1195 47.7 418 1214 53.8 403 1194 51.8 Alloy 24 EDM cutting 440 1166 31.0 443 1167 32.0 455 1176 31.0 Punched 496 696 5.0 463 688 5.0 440 684 4.0 Punched and annealed 559 1100 22.3 581 1113 22.0 561 1100 22.3 Alloy 25 EDM cutting 474 1183 15.8 470 1204 17.0 485 1223 17.4 Punched 503 589 2.1 517 579 0.8 497 583 2.1 Punched and annealed 457 1143 15.4 477 1159 14.6 423 1178 16.3 Alloy 26 EDM cutting 735 1133 20.8 742 1109 19.0 Punched 722 898 3.4 747 894 2.9 764 894 3.1 Punched and annealed 715 1112 18.8 713 1098 17.8 709 931 10.0 Alloy 27 EDM cutting 537 1329 19.3 513 1323 21.4 480 1341 20.8 Punched 563 624 4.3 568 614 3.3 539 637 4.3 Punched and annealed 505 1324 19.7 514 1325 20.0 539 1325 19.4 Alloy 29 EDM cutting 460 1209 54.7 441 1199 54.1 475 1216 52.9 Punched 489 828 15.4 486 811 14.6 499 813 14.8 Punched and annealed 410 1204 53.9 410 1220 53.2 408 1214 52.3 Alloy 32 EDM cutting 431 1196 50.6 437 1186 52.0 420 1172 54.7 Punched 471 826 19.9 452 828 19.7 482 854 19.7 Punched and annealed 406 1169 58.1 403 1170 51.4 405 1176 57.7

TABLE 32 Summary of Tensile Properties; Loss (E2 / E1) and Gain (E3 / E1)

alloy Loss of tensile elongation
(E2 / E1) Gain of tensile elongation
(E3 / E1) at least maximum at least maximum Alloy 1 0.14 0.21 0.88 1.06 Alloy 2 0.23 0.31 1.00 1.09 Alloy 9 0.41 0.56 0.93 1.13 Alloy 11 0.38 0.45 0.84 1.03 Alloy 12 0.42 0.57 0.83 1.21 Alloy 13 0.11 0.16 0.95 1.07 Alloy 14 0.28 0.31 0.84 0.99 Alloy 18 0.06 0.08 0.94 1.07 Alloy 21 0.20 0.31 0.83 1.10 Alloy 24 0.13 0.16 0.69 0.72 Alloy 25 0.05 0.13 0.89 1.03 Alloy 26 0.14 0.18 0.48 0.99 Alloy 27 0.15 0.22 0.91 1.04 Alloy 29 0.27 0.29 0.97 1.02 Alloy 32 0.36 0.39 0.94 1.15

Punching of the tensile test results in edge damage and degrades the tensile properties of the material. Plastic deformation of the sheet alloys herein during punching leads to structural transformation into reduced high malleability, high strength nanomodal structures (structure # 5, FIG. 1B), thereby premature cracking at edges and relatively low properties (e.g., , Reduction in elongation and tensile strength). This example shows that due to the inherent structural reversibility, the edge damages of the alloys listed in Table 2 are substantially recoverable by annealing, and recrystallized modal structures with full or partial recovery of properties depending on the alloy's chemical properties and process. It returns to formation of (Structure # 4, FIG. 1B). For example, as illustrated by alloy 1, the generation of punched, sheared and sheared edges decreased from an average of about 1310 MPa (EDM cut sample without sheared / damaged edges) to an average value of 678 MPa ( Decrease in the range of 45 to 50%). Upon annealing, the tensile strength returns to an average value of about 1308 MPa, which is in the range of at least 95% of the original value of 1310 MPa. Similarly, the tensile elongation is initially about 47.1% on average, reduced to an average value of 8.1% (a reduction of up to about 80 to 85%), and upon undergoing annealing and treatment shown by mechanism # 3 in FIG. Recovers to an average of 46.1% (more than 90% of elongation at 47.1%).

Example # 8 Temperature Effect on Recovery and Recrystallization

Slabs 50 mm thick from Alloy 1 were laboratory cast, hot rolled to a thickness of 2 mm in the laboratory, and cold rolled to a reduction of about 40%. Tensile specimens of ASTM E8 shape were prepared by wire EDM cutting from the cold rolled sheet. Some of the tensile specimens were annealed for 10 minutes at different temperatures in the range of 450-850 ° C. and then air cooled. Tensile properties were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Tensile test results are shown in FIG. 40 showing the transition of deformation behavior with annealing temperature. During the cold rolling process, dynamic nanophase reinforcement (mechanism # 2, FIG. 1A) or nanophase refining and reinforcement (mechanism # 4, FIG. 1B) occurs, which results from austenite to ferrite when the yield strength exceeds an increasing strain. And continuous transformation of nanoscale to hexagonal phase. Simultaneously with this transformation, deformation by a dislocation mechanism occurs in the matrix grains before and after transformation. The result is a change from the nanomodal structure (structure # 2, FIG. 1A) to the high strength nanomodal structure (structure # 3, FIG. 1A) or the recrystallized modal structure (structure # 4, FIG. 1B). # 5, a change to FIG. 1B occurs. Changes in structure and properties that occur during cold deformation can be reversed to varying degrees by annealing depending on the annealing parameters as shown in the tensile curve of FIG. 40. In FIG. 40B, the corresponding yield strength from the tensile curve is provided as a function of heat treatment temperature. The yield strength in the unannealed state after cold rolling is measured at 1141 MPa. As shown, depending on the annealing method of the material, which may include partial and complete recovery and partial and complete recrystallization, the yield strength is largely from 1372 MPa at annealing at 500 ° C. to 458 MPa at annealing at 850 ° C. Can be changed.

TEM irradiation was performed on selected samples annealed at different temperatures to account for the recovery of microstructures with tensile properties during annealing. For comparison, cold rolled sheets are included herein as reference. A laboratory cast alloy 1 slab of 50 mm thickness was used, which was hot rolled at 1250 ° C. in two steps of 80.8% and 78.3% to a thickness of about 2 mm, then cold rolled by 37% to 1.2 mm The sheet of thickness was produced. The cold rolled sheet was annealed at 450 ° C., 600 ° C., 650 ° C. and 700 ° C. for 10 minutes each. 41 shows the microstructure of Alloy 1 sample as cold rolled. It can be seen that after cold rolling a typical high strength nanomodal structure is formed, where a high density of dislocations is created with the presence of strong tissue. Annealing at 450 ° C. for 10 minutes resulted in the formation of recrystallization and high strength nanomodal structures, as the microstructures remained similar to those of cold rolled structures, and the rolling structure remained unchanged (FIG. 42). Does not cause When the cold rolled sample was annealed at 600 ° C. for 10 minutes, the TEM analysis showed very small isolated grains that signaled the onset of recrystallization. As shown in FIG. 43, isolated grains on the order of 100 nm are produced after annealing, and there is also a region of the modified structure with a potential network. Annealing at 650 ° C. for 10 minutes shows larger recrystallized grains, indicating the progress of recrystallization. Although the fraction of the deformed regions is reduced, the deformed structure continues to be seen as shown in FIG. 44. Annealing at 700 ° C. for 10 minutes shows larger and clear recrystallized grains as indicated in FIG. 45. Selected electron diffraction shows that these recrystallized grains are austenite phases. The area of the modified structure is small compared to the sample annealed to lower temperature. Examination of the entire sample suggests that about 10% to 20% of the area is occupied by the modified structure. The progress of recrystallization, as revealed by the TEM of the samples annealed at lower to higher temperatures, corresponds well to the change in tensile properties shown in FIG. 40. These low temperature (eg, less than 600 ° C.) annealed samples maintain predominantly high strength nanomodal structures leading to reduced malleability. The recrystallized sample (eg 700 ° C.) recovers most of the elongation compared to the sample completely recrystallized at 850 ° C. Annealing between these temperatures partially restores malleability.

One reason behind the difference in recovery and transition of deformation behavior is illustrated by the typical TTT diagram of FIG. 46. As mentioned above, very fine / nanoscale grains of ferrite formed during cold working recrystallize into austenite during annealing, and a fraction of a portion of the nano precipitates is redissolved. At the same time, the effects of strain hardening are removed along with dislocation networks, entanglements, twin interfaces, and incineration boundaries that are dissipated by various known mechanisms. As shown by the heating curve (A) of the typical temperature, time transformation (TTT) diagram of FIG. 46, at low temperatures (especially below 650 ° C. for Alloy 1), only recovery can occur without recrystallization (ie, recovery is Is the basis for the reduction of dislocation density).

In other words, in the broad context of the present invention, the effects of shear and the formation of sheared edges, and the negative effects on the mechanical properties associated therewith, are at least partially recovered at temperatures between 450 ° C. and 650 ° C. as shown in FIG. 46. Can be. In addition, at 650 ° C. to less than the Tm of the alloy, recrystallization may occur, which also contributes to the recovery of lost mechanical strength due to the formation of sheared edges.

Thus, this example shows that a simultaneous process involving dynamic strain hardening and phase transformation occurs through inherent mechanisms # 2 or # 3 (FIG. 1A) with dislocation-based mechanisms in deformation during cold rolling. Upon heating, the microstructure can be inverted into a recrystallized modal structure (structure # 4, FIG. 1B). However, at low temperatures, this reversal process may not occur if only potential recovery occurs. Thus, due to the inherent mechanism of the alloys of Table 2, a positive external heat treatment can be used to repair edge damage from punching / stamping.

Example # 9 Temperature Effect of Punched Edge Recovery

Slabs with a thickness of 50 mm were laboratory cast with the selected alloys listed in Table 33 according to the atomic ratios provided in Table 2, and were subjected to laboratory processing of hot rolling, cold rolling and 10 at 850 ° C. as described in the text portion of the present application. Annealed for minutes. Sheets obtained from each alloy with a final thickness of 1.2 mm and recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate recovery of punched edge damage after annealing as a function of temperature.

Tensile specimens of ASTM E8 shape were prepared by punching. A portion of the tensile specimen punched from the selected alloy was air cooled after recovery annealing for 10 minutes at different temperatures in the range of 450-850 ° C. Tensile properties were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer.

Tensile test results are shown in Table 32 and FIG. 47. As can be seen, complete or near complete recovery of properties has been achieved after annealing at temperatures above 650 ° C., which means that after punching the structure is completely or almost completely recrystallized at the damaged edges (ie, from structure # 5 to structure # 4 of FIG. 1B). Structural change). For example, the level of recrystallization at the damaged edge is considered to be at least 90% when the annealing temperature ranges from 650 ° C. to T _m . Lower annealing temperatures (eg, below 650 ° C.) do not cause complete recrystallization and cause partial recovery (ie, reduction of dislocation density) as described in Example # 8 and shown in FIG. 6.

Changes in the microstructure at the shear edges of Alloy 1 as a result of punching and annealing at different temperatures were examined by SEM. Sectional samples were cut from ASTM E8 punched tensile specimens near the sheared edges after annealing at 650 ° C. and 700 ° C. in the punched intact state as shown in FIG. 48.

For SEM irradiation, the cross section of the sample was polished with a small grit SiC abrasive paper and then gradually polished with a diamond media paste of up to 1 μm. Final polishing was performed using a 0.02 μm grit SiO ₂ solution. Microstructure was examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.

49 shows a backscatter SEM image of the microstructure of the edges as punched. In contrast to the recrystallized microstructure of the area away from the shear influence area, it can be seen that the microstructure within the shear influence area (a triangle with white contrast close to the edge) has been deformed and transformed. Similar to tensile deformation, the deformation of the shear influence area caused by punching produces a refined high strength nanomodal structure (structure # 5, FIG. 1B) through nanophase refinement and reinforcement mechanisms. However, the annealing restores the tensile properties of the punched ASTM E8 specimen, which is related to the change in the microstructure of the shear affected area during the annealing. 50 shows the microstructure of the sample annealed at 650 ° C. for 10 minutes. Compared to the sample as punched out, the shear influence area becomes smaller with less contrast, suggesting that the microstructure of the shear influence area evolves into the sample centered microstructure. High magnification SEM images show that some very small grains nucleate, but recrystallization does not occur on a large scale throughout the shear affected zone. This recrystallization is likely to be in the early stages when most of the dislocation has disappeared. Although the structure was not completely recrystallized, the tensile properties were substantially restored (Table 32 and Figure 47A). Annealing at 700 ° C. for 10 minutes results in complete recrystallization of the shear affected zone. As shown in FIG. 51, the contrast of the shear influence region is significantly reduced. The high magnification image shows that equiaxed grains with clear grain boundaries are formed in the shear influence region, indicating complete recrystallization. Grain size is smaller than that of the center of the sample. Note that the central grain was produced from recrystallization after annealing for 10 minutes at 850 ° C. before punching the specimen. With the shear affected zone completely recrystallized, the tensile properties are fully recovered, as shown in Table 32 and FIG. 47A.

Punching tensile specimens causes edge damage that degrades the tensile properties of the material. Plastic deformation of the sheet alloys herein during punching leads to structural transformations into reduced high malleability, high strength nanomodal structures (structure # 5, FIG. 1B) leading to premature cracking at the edges. This example demonstrates that edge damage is partially / completely recoverable by different annealing over a wide range of industrial temperatures.

TABLE 33 Tensile Properties after Punching and Annealing at Different Temperatures

alloy Annealing temperature
(℃) Yield strength
(MPa) Tensile strength
(MPa) Tensile elongation
(%) Alloy 1 AS Punched 494 798 12.6 487 829 14.3 474 792 15.3 450 481 937 21.5 469 934 20.9 485 852 19.3 600 464 1055 27.3 472 1103 30.5 453 984 23.7 650 442 1281 51.5 454 1270 45.4 445 1264 51.1 700 436 1255 50.1 442 1277 52.1 462 1298 51.6 850 407 1248 52.0 406 1260 47.8 412 1258 48.5 Alloy 9 AS Punched 508 1018 29.2 507 1007 28.6 490 945 23.3 600 461 992 28.5 462 942 24.8 471 968 25.6 650 460 1055 33.0 470 1166 48.3 473 1177 49.3 700 457 1208 57.5 455 1169 50.3 454 1171 61.6 850 411 1166 59.0 409 1174 52.7 418 1181 55.6 Alloy 12 AS Punched 521 954 27.1 468 978 30.7 506 975 31.2 600 462 1067 44.9 446 1013 41.3 471 1053 41.1 650 452 1093 61.5 449 1126 57.8 505 1123 55.4 700 480 1112 59.6 460 1117 61.8 468 1096 61.5 850 419 1086 65.7 423 1085 63.0 415 1100 53.8

Example # 10 Effect of Punching Speed on Punched Edge Characteristics Reversibility

Slabs with a thickness of 50 mm were laboratory cast with selected alloys listed in Table 34 according to the atomic ratios given in Table 2, and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes as described herein. It was. Sheets obtained from each alloy with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate edge damage recovery as a function of punching speed.

Tensile specimens of ASTM E8 shape were prepared by punching at three different speeds: 28 mm / sec, 114 mm / sec, and 228 mm / sec. Wire EDM cut specimens of the same material were used for reference. A portion of the tensile specimen punched from the selected alloy was air cooled after recovery annealing at 850 ° C. for 10 minutes. Tensile properties were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grips were fixed and the upper grips moved upwards at a rate of 0.012 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Tensile test results are reported in Table 34, and tensile properties as a function of punch speed of the selected alloy are shown in FIG. Tensile properties were found to be significantly reduced in the punched samples compared to the tensile properties of wire EDM cuts. An increase in punching speed from 28 mm / sec to 228 mm / sec improves the properties of all three selected alloys. Local heating during punching holes or shearing edges is known to increase with increasing punching speed and may be a factor in recovering edge damage of specimens punched at higher speeds. Note that heat alone does not cause edge damage recovery but is made possible by the reaction of the material to the heat produced. This difference in the reaction of the alloys included in Table 2 herein for commercial steel samples is clearly illustrated in Examples 15 and 17.

TABLE 34 Tensile Properties of Specimens Punched and EDM Cut at Different Speeds

alloy Sample preparation method Yield strength
(MPa) The tensile strength
(MPa) Tensile elongation
(%) Alloy 1 EDM 459 1255 51.2 443 1271 46.4 441 1248 52.7 453 1251 55.0 467 1259 51.3 Punched 228 mm / sec 474 952 21.8 498 941 21.6 493 956 21.6 114 mm / sec punched 494 798 13.4 487 829 15.1 474 792 14.1 Punched 28 mm / sec 464 770 12.8 479 797 13.7 465 755 12.1 Alloy 9 EDM 468 1166 56.1 480 1177 52.4 475 1169 56.9 Punched 228 mm / sec 500 1067 35.1 493 999 28.8 470 1042 31.8 114 mm / sec punched 508 1018 29.2 507 1007 28.6 490 945 23.3 Punched 28 mm / sec 473 851 19.7 472 841 16.4 494 846 18.9 Alloy 12 EDM 481 1094 54.4 479 1128 64.7 495 1126 62.4 Punched 228 mm / sec 495 1124 53.8 484 1123 53.0 114 mm / sec punched 521 954 27.1 468 978 30.7 506 975 31.2 Punched 28 mm / sec 488 912 23.6 472 900 21.7 507 928 22.9

This example demonstrates that punch rates can have a significant impact on the obtained tensile properties of the steel alloys herein. Local heat generation during punching can be a factor in structural recovery near the edges leading to improved characteristic properties.

Example # 11 Edge structure transformation during hole punching and expansion

Slabs with a thickness of 50 mm were lab cast from Alloy 1 and annealed for 10 minutes at 850 ° C. by hot rolling, cold rolling and laboratory as described herein. The resulting sheet with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) was used for the hole expansion rate (HER) test.

Specimens measuring 89 × 89 mm were wire EDM cut from the sheet. A 10 mm diameter hole was cut in the middle of the specimen using two methods: punching and drilling by edge milling. Hole punching was performed using a constant speed of 0.25 mm / sec and 16% clearance on an Instron Model 5985 general purpose test system. The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation. Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The results of the HER test are shown in FIG. 53, demonstrating significantly lower values (5.1% HER vs. 73.6% HER, respectively) compared to those by milling holes formed by punching holes. Samples were cut from two samples tested as shown in FIG. 54 for SEM analysis and microhardness measurement.

The microhardness of Alloy 1 was measured at all relevant stages of the hole expansion process. Microhardness measurements were made along the cross section of the sheet sample in the annealed state (prior to punching and HER test), punched state, and HER tested state. The microhardness of the cold rolled sheet of alloy 1 was also measured for reference. The measurement profile started at a distance of 80 microns from the edge of the sample, with further measurements taken every 120 microns up to 10 times. After that point, additional measurements were taken every 500 microns to determine a total sample length of at least 5 mm. A schematic of the microhardness measurement position of the HER tested sample is shown in FIG. 55. SEM images of the punched and HER tested samples after microhardness measurements are shown in FIG. 56.

As shown in FIG. 57, the punching process creates a strained region of about 500 microns directly adjacent to the punched edge, with the material closest to the punched edge completely or almost completely transformed, which is next to the punched edge. This is evidenced by the hardness approach observed in fully transformed 40% cold rolled material. The microhardness profile for each sample is shown in FIG. 58. As can be seen, when milled, the microhardness gradually increased towards the hole edge, while the increase in the microhardness in the case of punched holes was observed in a very narrow area proximate to the hole edge. As shown in FIG. 58, the TEM samples of both cases were cut at the same distance.

To prepare the TEM specimens, the HER test samples were first cut by wire EDM, and pieces with hole edge portions were thinned by grinding with pads with reduced grit size. Further thinning up to about 60 μm thick was performed by polishing with diamond suspensions of 9 μm, 3 μm, and 1 μm, respectively. A 3 mm diameter disc was punched from near the edge of the hole of the foil and final polishing was completed by electropolishing using a twin jet polisher. The chemical solution used was 30% nitric acid mixed in methanol base. For insufficient thin regions for TEM observations, TEM specimens can be ion milled using the Gatan Precision Ion Polishing System (PIPS). Ion milling is typically performed at 4.5 keV and the incline angle is reduced from 4 ° to 2 ° to exploit thin areas. TEM irradiation was performed using a JEOL 2100 high resolution microscope operated at 200 kV. Since the location for the TEM irradiation is in the center of the disk, the observed microstructure is about 1.5 mm from the edge of the hole.

The initial microstructure of the alloy 1 sheet prior to testing is shown in FIG. 69 showing the recrystallized modal structure (structure # 4, FIG. 1B). FIG. 60A shows a TEM micrograph of the microstructure of a HER test sample with punched holes (HER = 5.1%) after testing in different areas at 1.5 mm from the hole edge. The recrystallized microstructure was found to remain in the sample (FIG. 60A) with a small area with mainly partially transformed “pockets” (FIG. 60B), which indicated that the limited volume of the sample (about 1500 μm depth) in the HER test. Related to the variant. As shown in FIG. 61, in a HER sample with milled holes (HER = 73.6%), the sample as indicated by a large amount of transformed “pockets” and high density dislocations (10 ⁸ to 10 ¹⁰ mm ⁻² ) There is a large amount of variation in it.

In order to analyze in more detail why HER performance is poor in samples with punched holes, Focused Ion Beam (FIB) technology was used to prepare TEM specimens at the edges of the punched holes. As shown in FIG. 62, the TEM specimen is cut at about 10 μm from the edge. To prepare a TEM specimen by FIB, a thin layer of platinum is deposited in that area to protect the specimen to be cut. The wedge shaped specimen is then cut out and lifted with a tungsten needle. Additional ion milling is performed to thin the specimen. Finally, the thinned specimen is transferred and welded to a copper grid for TEM observation. FIG. 63 shows the microstructure of Alloy 1 sheet at a distance of about 10 microns from the punched hole edge significantly refined and transformed relative to the microstructure of Alloy 1 sheet before punching. This is because punching causes severe deformation at the hole edges, resulting in nanophase refinement and reinforcement (mechanism # 4, FIG. 1B), resulting in a refined high strength nanomodal structure (structure # 5, FIG. 1B) in the region proximate the punched hole edges. Suggests that it leads to formation. This structure has a relatively lower malleability compared to the recrystallized modal structure Table 1, resulting in premature cracking at the edges and low HER values. This example shows a high-strength nanomodal structure (structure # 5, micronized from modal structures (structure # 4, FIG. 1b) recrystallized through nanophase micronization and reinforcement (mechanism # 4, FIG. 1B), in which the alloys of Table 2 were specified. It demonstrates the inherent ability to metamorphosis to Figure 1b). Structural transformations resulting from deformation at the hole edges during punching appear to be similar in nature to those observed during deformation and tensile test deformations during cold rolling deformation.

Example # 12 HER test results with or without annealing

Slabs with a thickness of 50 mm were laboratory cast with selected alloys listed in Table 35 according to the atomic ratios provided in Table 2, and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes as described herein. It was. The resulting sheet with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) was used for the hole expansion rate (HER) test.

89 x 89 mm specimens were wire EDM cut from a larger portion of the sheet. A 10 mm diameter hole was formed in the center of the specimen by punching using a constant speed of 0.25 mm / s at a punch spacing of 6% on an Instron Model 5985 universal test system. Half of the prepared specimens with punched holes were individually wrapped in stainless steel foil and annealed at 850 ° C. for 10 minutes before the HER test. The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

Table 35 shows the results of measuring hole slide rates for specimens annealed and unannealed after hole punching. As shown in FIGS. 64, 65, 66, 67, and 68 for Alloy 1, Alloy 9, Alloy 12, Alloy 13, and Alloy 17, respectively, measured in the case of punched holes that have been annealed. Hole expansion rates are greater than unannealed punched holes. An increase in hole expansion rate in the annealing state for the alloys specified herein leads to an increase in the actual HER of about 25% to 90%.

Table 35. Hole Expansion Rate Results for Selected Alloys with and Without Annealing

material condition Punch interval
(%) Measured Hole Expansion Rate (%) Average hole expansion rate
(%) Alloy 1 No annealing 16 3.00 3.20 3.90 2.70 Annealing 16 105.89 93.10 81.32 92.11 Alloy 9 No annealing 16 3.09 3.19 3.19 3.29 Annealing 16 78.52 87.84 97.60 87.40 Alloy 12 No annealing 16 4.61 4.91 5.21 Annealing 16 69.11 77.60 83.60 80.08 Alloy 13 No annealing 16 1.70 1.53 1.40 1.50 Annealing 16 32.37 31.12 29.00 32.00 Alloy 17 No annealing 16 12.89 21.46 28.70 22.80 Annealing 16 104.21 103.74 80.42 126.58

This example shows that the edge formability demonstrated during the HER test can lead to poor results due to edge damage during punching operations as a result of the alloy's inherent mechanisms listed in Table 2. Fully post-treated alloys exhibit very high tensile legends as shown in Tables 6-10 by engaging resistance to necking leading to very high strain hardening and near fracture. Thus, the material is highly resistant to the worst failures, but artificial worst failures occur near the punched edge during punching. Due to the inherent reversibility of the specified mechanisms, these harmful edge damage as a result of nanophase refinement and enhancement (mechanism # 3, FIG. 1A) and structural transformation can be reversed by annealing to obtain high HER results. Thus, when annealing is performed after punching the hole, a high hole expansion ratio value can be obtained, and an excellent combination of tensile properties and associated bulk formability can be maintained.

In addition, the alloy of the present disclosure via a processing path for providing such an alloy in the form of structure # 4 (a recrystallized modal structure), in the case of a hole formed by shearing and comprising a sheared edge, has a first hole expansion. It can be understood that the ratio HER ₁ , and heating this alloy, has a second hole expansion rate HER ₂ , where HER ₂ > HER ₁ .

More specifically, the alloys of the present disclosure via a processing path for providing such alloys with structure # 4 (recrystallized modal structure) are not primarily dependent on shear compared to punching holes (ie It can be understood that, for holes installed in the alloy via a method (i.e., waterjet cutting, laser cutting, wire-edm, milling, etc.), the value represents a first hole expansion ratio (HER ₁ ) in the range of 0 to 130%. However, if the same alloy contains holes formed by shearing, a second hole expansion ratio HER ₂ is observed where HER ₂ = (0.01 to 0.30) (HER ₁ ). , HER ₂ is observed to recover to HER ₃ = (0.60 to 1.0) HER ₁ .

Example # 13 Effect of Edge Conditions on Alloy Properties

Slabs with a thickness of 50 mm were lab cast into alloy 1 according to the atomic ratios provided in Table 2, and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes as described herein. Sheets obtained from Alloy 1 with a final thickness of 1.2 mm and recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate the effect of edge conditions on the tensile and hole expansion properties of Alloy 1.

Tensile specimens of ASTM E8 shape were fabricated using two methods, punching and wire EDM cutting. Punched tensile specimens were prepared using a commercial press. Some of the punched tensile specimens were heat treated at 850 ° C. for 10 minutes to produce samples with annealed edge conditions after punching.

Tensile properties of ASTM E8 specimens were measured on an Instron 5984 mechanical test frame using Instron's Bluehill control software. All tests were performed at room temperature, the lower grip was fixed, and the upper grip was first set to move upwards at a rate of 0.025 mm / sec with a target of 0.5% elongation and after that at 0.125 mm / sec. Strain data was collected using Instron's Advanced Video Extensometer. Table 36 shows the tensile properties of Alloy 1 with punched, EDM cut, and annealed edge conditions after punching. 69 shows the tensile properties of Alloy 1 with different edge conditions.

Table 36 Tensile Properties of Alloy 1 with Different Edge Conditions

Edge condition Tensile Elongation (%) Tensile strength (MPa) Punched 12.6 798 14.3 829 15.3 792 EDM cutting 50.5 1252 51.2 1255 52.7 1248 55.0 1251 51.3 1259 50.5 1265 Annealed after punched 52.0 1248 47.8 1260 48.5 1258

Specimens for hole expansion rate test with dimensions of 89 × 89 mm were wire EDM cut from the sheet. 10 mm diameter holes were prepared in two ways: punching and wire EDM cutting. Punched holes of 10 mm diameter were formed by punching at 0.25 mm / sec using 16% punch spacing and planar punch profile shapes on an Instron 5985 universal test system. A portion of the punched sample for hole expansion rate test was annealed by heat treatment at 850 ° C. for 10 minutes after punching.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The hole expansion rate test results are shown in Table 37. The average hole expansion rate values for each edge condition are also shown. The average hole expansion rate for each edge condition is shown in FIG. For samples with annealed edge conditions after EDM cutting and punching, the edge formability (ie HER response) is excellent, whereas samples with holes in punched edge conditions have significantly lower edge formability. Able to know.

Table 37. Hole Expansion Rate of Alloy 1 with Different Edge Conditions

Edge condition Measured Hole Expansion Rate (%) Average Hole Expansion Rate (%) Punched 3.00 3.20 3.90 2.70 EDM cutting 92.88 82.43 67.94 86.47 Annealed after punched 105.90 93.10 81.30 92.10

This example demonstrates that the edge conditions of Alloy 1 have a marked effect on tensile properties and edge formability (ie, HER response). Tensile samples tested with punched edge conditions have attenuated properties compared to both wire EDM cut and annealed after punch. Samples with punched edge conditions have an average 3.20% hole expansion rate while EDM cut and punched and annealed edge conditions have 82.43% and 93.10% hole expansion rates, respectively. Comparison of edge conditions also demonstrates that damage related to edge formation (ie, through punching) has a significant impact on the edge formability of the alloys herein.

Example # 14 HER Results as a Function of Hole Punching Speed

Slabs having a thickness of 50 mm were laboratory cast with selected alloys listed in Table 38 according to the atomic ratios given in Table 2, and hot rolled, cold rolled and annealed for 10 minutes at 850 ° C. as described herein. It was. Sheets obtained from each alloy with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate the effect of hole punching speed on HER results.

Specimens measuring 89 × 89 mm were wire EDM cut from the sheet. Holes with a diameter of 10 mm were punched at different speeds on two different machines, but all specimens were punched to 16% punch spacing and the same punch profile shape. Low speed (0.25 mm / sec, 8 mm / sec) punched holes were punched using the Instron 5985 universal test system, and high speed (28 mm / sec, 114 mm / sec, 228 mm / sec) punched holes were commercial punches. Punched on press. All holes were punched using a flat punch shape.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The hole expansion rate values for the test are shown in Table 37. Average hole expansion values are shown for the alloys tested at each speed and 16% punch interval. Average hole expansion rates as a function of punch speed are shown in FIGS. 71, 72 and 73 for Alloy 1, Alloy 9, and Alloy 12, respectively. It can be seen that as the punch speed increases, all alloys tested have a positive edge formability response as evidenced by the increase in hole expansion rate. The reason for this increase is considered to be related to the following effects. Higher punch rates are expected to increase the amount of heat generated at the sheared edges, which local temperature spikes can cause an annealing effect (ie in situ annealing). Alternatively, increasing the punch speed may reduce the amount of material transformed from the recrystallized modal structure (ie, structure # 4 in FIG. 1B) to the refined high strength nanomodal structure (ie, structure # 5 in FIG. 1B). Can be. At the same time, due to temperature spikes that enable local recrystallization (ie, mechanism # 3 of FIG. 1B), the amount of micronized high-strength nanomodal structure (ie, structure # 5 in FIG. 1B) may be reduced.

Table 38. Hole Expansion Rate at Different Punch Speeds

material Punch speed
(mm / sec) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) Alloy 1 0.25 3.00 3.20 0.25 3.90 0.25 2.70 8 4.49 3.82 8 3.49 8 3.49 28 8.18 7.74 28 8.08 28 6.97 114 17.03 17.53 114 19.62 114 15.94 228 20.44 21.70 228 21.24 228 23.41 Alloy 9 0.25 3.09 3.19 0.25 3.19 0.25 3.29 8 6.80 6.93 8 7.39 8 6.59 28 21.04 19.11 28 17.35 28 18.94 114 24.80 24.29 114 19.74 114 28.34 228 26.00 30.57 228 35.16 228 30.55 Alloy 12 0.25 4.61 4.91 0.25 5.21 8 7.62 11.28 8 14.61 8 11.62 28 29.38 31.59 28 33.70 28 31.70 114 40.08 45.50 114 48.11 114 48.31 228 50.00 49.36 228 40.56 228 57.51

This example demonstrates the dependence of the edge formability on the punching speed measured by the hole expansion rate. As the punch speed increases, the hole expansion rate for the alloys tested generally increases. As the punching speed increases, the properties of the edges change to improve edge formability (ie, HER response). At punching speeds greater than measured, edge formability is expected to continue to improve towards much larger hole expansion rate values.

Example # 15 HER of DP980 as a Function of Hole Punching Speed

The hole expansion rate test was performed by purchasing a commercially produced and processed dual phase 980 steel. All specimens were tested as received (commercially processed).

Specimens measuring 89 × 89 mm were wire EDM cut from the sheet. Holes with a diameter of 10 mm were punched at different speeds on two different machines, but all specimens were punched with a punch punch of 16% and the same punch profile shape using a commercial punch press. Low speed (0.25 mm / sec) punched holes were punched using an Instron 5985 universal test system, and high speed (28 mm / sec, 114 mm / sec, 228 mm / sec) punched holes were punched on a commercial punch press. All holes were punched using a flat punch shape.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

Hole expansion test values are shown in Table 39. Average hole expansion values for each punch speed are also shown for commercial dual phase 980 materials at 16% punch spacing. The average hole expansion value in FIG. 74 is plotted as a function of punching speed for commercial dual phase 980 steels.

Table 39. Hole expansion rate for dual phase 980 steels at different punch rates

material Punch speed
(mm / sec) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) Commercial Dual Phase 980 0.25 23.55 22.45 0.25 20.96 0.25 22.85 28 18.95 18.26 28 17.63 28 18.21 114 17.40 20.09 114 23.66 114 19.22 228 27.21 23.83 228 24.30 228 19.98

This example shows that in dual phase 980 steels, edge performance effects based on punch speed cannot be measured. For all punch rates measured in the dual phase 980 steel, the edge performance (ie HER response) remains consistent within the range of 21% ± 3%, as presented herein in FIGS. 1A and 1B. Since there is no unique structure and mechanism, the edge performance of conventional AHSS is not improved by the punch speed as expected.

Example # 16: HER Results as a Function of Punch Design

Slabs with a thickness of 50 mm were laboratory cast into Alloy 1, Alloy 9 and Alloy 12 according to the atomic ratios provided in Table 2, and were subjected to laboratory treatment for hot rolling, cold rolling and 10 minutes at 850 ° C. as described herein. Annealed. Sheets obtained from each alloy with a final thickness of 1.2 mm and a recrystallized modal structure (structure # 4, FIG. 1B) were used to demonstrate the effect of hole punching speed on HER results.

89 x 89 mm specimens were wire EDM cut from larger portions. Punching a 10 mm diameter hole in the center of the specimen with three different speeds (28 mm / sec, 114 mm / sec, and 228 mm / sec), 16% punch spacing, and four punch profile shapes using a commercial punch press It was. These punch shapes used were flat, 6 ° tapered, 7 ° conical, and conical planes. A schematic of the punch shapes of 6 ° tapered, 7 ° conical, and conical planes is shown in FIG.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

Hole expansion rate data are included in Tables 40, 41, and 42, respectively for Alloy 1, Alloy 9, and Alloy 12 at four punch shapes and two punch rates. Average hole expansion values for Alloy 1, Alloy 9, and Alloy 12 are shown in FIGS. 76, 77, and 78, respectively. For all alloys tested, the 7 ° conical punch shape achieved an expansion rate equal to the maximum or maximum hole expansion rate compared to all other punch shapes. Increasing the punch rate also improves edge formability (ie HER response for all punch shapes). If the punching speed is increased at different punch shapes, the alloys of the present disclosure will provide for local heating of the edges at these higher relative punch rates, It is believed that there may be a trigger mechanism (# 3), and some formation of structure (# 4), so that some recrystallization (mechanism # 3) may be experienced.

Table 40. Hole Expansion Rate of Alloy 1 at Different Punch Shapes

Punch shape Punch speed
(mm / sec) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) plane 28 8.18 7.74 plane 28 8.08 plane 28 6.97 plane 114 17.03 17.53 plane 114 19.62 plane 114 15.94 plane 228 20.44 21.70 plane 228 21.24 plane 228 23.41 6 ° taper 28 7.87 8.32 6 ° taper 28 8.77 6 ° taper 114 19.84 18.48 6 ° taper 114 16.55 6 ° taper 114 19.04 7 ° conical 28 8.37 10.56 7 ° conical 28 12.05 7 ° conical 28 11.25 7 ° conical 114 23.41 22.85 7 ° conical 114 21.14 7 ° conical 114 24.00 7 ° conical 228 21.71 21.37 7 ° conical 228 19.50 7 ° conical 228 22.91 Conical plane 28 8.47 11.95 Conical plane 28 13.25 Conical plane 28 14.14 Conical plane 114 20.42 19.75 Conical plane 114 19.22 Conical plane 114 19.62 Conical plane 228 24.13 22.39 Conical plane 228 23.31 Conical plane 228 19.72

Table 41. Hole Expansion Rate of Alloy 9 at Different Punch Shapes

Punch shape Punch speed
(mm / sec) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) plane 28 21.04 19.11 plane 28 17.35 plane 28 18.94 plane 114 24.80 24.29 plane 114 19.74 plane 114 28.34 plane 228 26.00 30.57 plane 228 35.16 plane 228 30.55 6 ° taper 28 17.35 19.36 6 ° taper 28 19.06 6 ° taper 28 21.66 6 ° taper 114 29.64 31.14 6 ° taper 114 32.14 6 ° taper 114 31.64 7 ° conical 28 22.63 24.05 7 ° conical 28 23.61 7 ° conical 28 25.92 7 ° conical 114 34.36 32.60 7 ° conical 114 31.67 7 ° conical 114 31.77 7 ° conical 228 36.28 36.44 7 ° conical 228 38.87 7 ° conical 228 34.16 Conical plane 28 27.72 25.59 Conical plane 28 24.63 Conical plane 28 24.43 Conical plane 114 30.28 32.64 Conical plane 114 32.87 Conical plane 114 34.76 Conical plane 228 32.90 35.45 Conical plane 228 37.45 Conical plane 228 35.99

Table 42. Hole Expansion Rate of Alloy 12 at Different Punch Shapes

Punch shape Punch speed
(mm / sec) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) plane 28 29.38 31.59 plane 28 33.70 plane 28 31.70 plane 114 40.08 45.50 plane 114 48.11 plane 114 48.31 plane 228 50.00 49.36 plane 228 40.56 plane 228 57.51 6 ° taper 28 29.91 30.67 6 ° taper 28 32.50 6 ° taper 28 29.61 6 ° taper 114 38.42 41.19 6 ° taper 114 44.37 6 ° taper 114 40.78 7 ° conical 28 34.90 33.76 7 ° conical 28 33.00 7 ° conical 28 33.37 7 ° conical 114 45.72 49.10 7 ° conical 114 49.30 7 ° conical 114 52.29 7 ° conical 228 58.90 54.36 7 ° conical 228 53.43 7 ° conical 228 50.75 Conical plane 28 37.15 34.43 Conical plane 28 31.47 Conical plane 28 34.66 Conical plane 114 45.76 46.36 Conical plane 114 45.96 Conical plane 114 47.36 Conical plane 228 57.51 54.11 Conical plane 228 53.48 Conical plane 228 51.34

This example shows that there is an impact of punch shape on edge formability for all alloys tested. For all alloys tested, the conical punch shape produces the maximum hole expansion rate, which means that modifying the punch shape from a planar punch to a conical punch shape reduces edge damage in the material due to punched edges, thereby improving edge formability. Prove that. The 7 ° conical punch shape shows the maximum edge formability increase compared to the planar punch shape, and the conical planar shape produces slightly lower hole expansion rates over most of the alloys tested. For Alloy 1, the effect of the punch shape is reduced with increasing punching speed, and the three tested shapes exhibit nearly equivalent edge formability as measured by the hole expansion rate (FIG. 79). It has been demonstrated that the punch shape greatly reduces residual damage from punching in the edges of the material with increasing punch speed, and as a result improves edge formability. Higher punch rates are expected to increase the amount of heat generated at the sheared edges, which local temperature spikes can cause an annealing effect (ie in situ annealing). Alternatively, increasing the punch speed may reduce the amount of material transformed from the recrystallized modal structure (ie, structure # 4 in FIG. 1B) to the refined high strength nanomodal structure (ie, structure # 5 in FIG. 1B). Can be. At the same time, due to temperature spikes that enable local recrystallization (ie, mechanism # 3 of FIG. 1B), the amount of micronized high-strength nanomodal structure (ie, structure # 5 in FIG. 1B) may be reduced.

Example # 17: Commercial steel grade HER as a function of hole punching speed

Hole expansion rate tests were performed on commercial steel grades 780, 980 and 1180. All specimens were tested as received sheets (commercially processed).

Specimens of 89 × 89 mm size were wire EDM cut from each grade of sheet. Holes with a 10 mm diameter were punched at different speeds on two different machines in the same punch profile shape using a commercial punch press. Low speed (0.25 mm / sec) punched holes were punched using an Instron 5985 universal test system, and high speed (28 mm / sec, 114 mm / sec, 228 mm / sec) punched holes were punched on a commercial punch press. Holes were punched using a flat punch shape.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The results of the hole expansion test are shown in Tables 43-45 and shown in FIG. 80. As can be seen, the hole expansion rate does not improve with increasing punching speed in all tested grades.

Table 43. Hole expansion rate of 780 steel grades at different punch speeds

Sample # Punch Speed (mm / sec) Percentage of Punch to Die Punch shape HER One 5 mm / s 12% plane 44.74 2 12% plane 39.42 3 12% plane 44.57 One 28 mm / sec 16% plane 35.22 2 16% plane 28.4 3 16% plane 36.38 One 114 mm / s 16% plane 31.58 2 16% plane 33.9 3 16% plane 22.29 One 228 mm / s 16% plane 31.08 2 16% plane 31.85 3 16% plane 31.31

Table 44 Hole expansion rates of 980 steel grades at different punch speeds

Sample # Punch Speed (mm / sec) % Spacing of punches to die Punch shape HER One 5 mm / s 12% plane 33.73 2 12% plane 35.02 One 28 mm / sec 16% plane 26.88 2 16% plane 26.44 3 16% plane 23.83 One 114 mm / s 16% plane 26.81 2 16% plane 30.56 3 16% plane 29.24 One 228 mm / s 16% plane 30.06 2 16% plane 30.98 3 16% plane 30.62

Table 45 Hole Expansion Rate of 1180 Steel Grades at Different Punch Speeds

Sample # Punch Speed (mm / sec) Percentage of Punch to Die Punch shape HER One 5 mm / s 12% plane 26.73 2 12% plane 32.9 3 12% plane 25.4 One 28 mm / sec 16% plane 35.32 2 16% plane 32.11 3 16% plane 36.37 One 114 mm / s 16% plane 35.15 2 16% plane 30.92 3 16% plane 32.27 One 228 mm / s 16% plane 27.25 2 16% plane 23.98 3 16% plane 31.18

This example shows that edge performance effects based on hole punching rates cannot be measured in commercial steel grades tested, since there is no inherent structure and mechanism presented herein as provided in FIGS. 1A and 1B. It is shown that the edge performance of conventional AHSS is not affected or improved by the punch speed as expected.

Example # 18: Correlation of Post Uniform Elongation to Hole Expansion Rate

Conventional steel materials have been found to show a strong correlation between the measured hole expansion rate and the post uniform elongation of the material. Post-uniform elongation of a material is defined as the difference between the total elongation of the sample in the tensile test and the uniform elongation at typical maximum tensile strength during the tensile test. Uniaxial tensile tests and hole expansion rate tests were completed for Alloy 1 and Alloy 9 on sheet materials about 1.2 mm thick for comparison with existing material correlations.

Slabs with a thickness of 50 mm were laboratory cast into Alloy 1 and Alloy 9 according to the atomic ratios given in Table 2, and hot rolled, cold rolled by laboratory treatment for 10 minutes at 850 ° C. as described in the text part of the present application. It was.

Tensile specimens of ASTM E8 shape were prepared by wire EDM. All samples were tested according to the standard test procedure described in the text section of this specification. The post uniform elongation was calculated using the average of the elongation and total elongation for each alloy. The average uniform elongation, average total elongation, and the calculated uniform elongation for Alloy 1 and Alloy 9 are provided in Table 46.

Specimens for hole expansion rate testing with dimensions of 89 × 89 mm were wire EDM cut from sheets of Alloy 1 and Alloy 9. 10 mm diameter holes were punched at 0.25 mm / sec on the Instron 5985 universal test system at 12% intervals. All holes were punched using a flat punch shape. These test parameters were chosen because they are commonly used by industry and academic experts for hole expansion rate testing.

The Hole Expansion Ratio (HER) test was performed on an SP-225 hydraulic press and consisted of slowly raising a conical punch that evenly extended the hole radially outward. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The measured hole expansion rate values for Alloy 1 and Alloy 9 are shown in Table 46.

Table 46. Uniaxial Tensile and Hole Expansion Data for Alloy 1 and Alloy 9

alloy Average
Uniform elongation Average
Total elongation Uniform elongation after
(ε _pul ) Hole expansion rate (%) (%) (%) (%) Alloy 1 47.19 49.29 2.10 2.30 Alloy 9 50.83 56.99 6.16 2.83

Commercial reference data are described by Paul S.K., J Mater Eng Perform 2014; 23: 3610.], Which is shown in Table 47. For commercial data, S.K. Paul predicts that the hole expansion rate of the material is proportional to 7.5 times the post-uniform elongation (see Equation 1).

HER = 7.5 (ε _pul ) Equation 1

TABLE 47 Reference data from Paul SK, J Mater Eng Perform 2014; 23: 3610.

Commercial steel grades Uniform elongation Total elongation Uniform elongation after
(ε _pul ) Hole expansion rate (%) (%) (%) (%) IF-Rephos 22 37.7 15.7 141.73 IF-Rephos 22.2 39.1 16.9 159.21 BH210 19.3 37.8 18.5 151.96 BH300 16.5 29 12.5 66.63 DP 500 18.9 27.5 8.6 55.97 DP600 16.01 23.51 7.5 38.03 TRIP 590 22.933 31.533 8.6 68.4 TRIP 600 19.3 27.3 8 39.98 TWIP940 64 66.4 2.4 39.1 HSLA 350 19.1 30 10.9 86.58 340 R 22.57 36.3 13.73 97.5

Post uniform elongation and hole expansion rates of Alloy 1 and Alloy 9 were obtained from commercial alloy data and S.K. It is plotted in FIG. 81 using Paul's expected correlation. Note that the data of Alloy 1 and Alloy 9 do not follow the expected correlation lines. This example does not follow the correlation of post uniform elongation and hole expansion rate for commercial steel grades for the steel alloys herein. Shows. The measured hole expansion rates of Alloy 1 and Alloy 9 are much smaller than predicted values based on the correlation of existing commercial steel grades, which, for example, inherent in structures and mechanisms as shown in FIGS. Influence is present in the steel alloys herein.

Example # 19 HER Performance as a Function of Hole Expansion Rate

Slabs with a thickness of 50 mm were lab cast into three selected alloys according to the atomic ratios provided in Table 2, and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes as described herein. Sheets from each alloy with recrystallized modal structures with a final thickness of 1.2 mm were used to demonstrate the effect of hole expansion rate on HER performance.

Specimens of 89 × 89 mm size were cut from the sheet with wire EDM. 10 mm diameter holes were punched on a commercial punch press at a constant speed of 228 mm / sec. All holes were punched out with a flat punch shape and spacing of punches for about 16% die.

Hole expansion rate (HER) testing was performed on an Interlaken Technologies SP-225 hydraulic press and consisted of raising a conical punch that evenly extended the hole radially outward. Four hole expansion speeds of 5, 25, 50, and 100 mm / min, ie conical ram movement speed, were used. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The hole expansion rate values for the test are shown in Table 48. Average hole expansion rate values are indicated for each speed, and the tested alloys show that the HER value increases as the hole expansion speed increases in all three alloys. The effect of the hole expansion rate is also demonstrated in FIGS. 82, 83, and 84 for Alloy 1, Alloy 9, and Alloy 12, respectively.

Table 48. Hole Expansion Rate of Selected Alloys at Different Expansion Rates

material Punch speed
(mm / sec) Hole expansion speed
(mm / min) Measured Hole Expansion Rate
(%) Average hole expansion rate
(%) Alloy 1 228 5 19.09 20.55 22.54 20.02 25 30.70 28.58 29.14 25.91 50 34.05 34.63 36.43 33.42 100 37.11 37.19 38.52 35.93 Alloy 9 228 5 34.06 34.15 34.07 34.31 25 32.87 40.77 45.46 43.98 50 38.39 44.17 39.71 54.42 100 48.01 49.50 55.27 45.23 Alloy 12 228 5 48.61 43.51 34.79 47.14 25 42.13 50.64 57.82 51.96 50 63.77 62.97 68.46 56.68 100 57.79 56.73 49.28 63.11

This example shows that the edge formability measured by the HER test, i.e., the deformation ability with relatively reduced cracking, can be influenced by the strain rate (i.e., hole expansion rate) of the hole edge. The alloy tested in this example shows a positive correlation between the hole expansion rate and the hole expansion rate where the measured hole expansion rate is relatively higher as the hole expansion rate is increased.

Thus, in the broad context of the present disclosure, by any edge forming method that causes deformation of the metal alloy when forming the edge (eg, by punching, shearing, piercing, drilling, cutting, cropping, stamping) It has been established that after the edges of the shape of are formed, increasing the speed of expanding these formed edges can expand further with the tendency of the relatively reduced cracking. Thus, the edges herein can include inner holes in metal sheets of alloys described herein or edges that form outer edges on such metal sheets. In addition, the edges herein can be formed by progressive die stamping operations, which are typically the basis for metalworking operations including punching, shearing, coining, and bending. Edges herein may be present in a vehicle, or more specifically in a vehicle frame, vehicle chassis, or part of a vehicle panel.

Edge expansion is understood herein to increase the length of such an edge with a corresponding change in edge thickness. This is confirmed by the above data in Table 48, which, in relation to the edges present in the holes, increases the rate of hole expansion (i.e., the edges in the holes are original when the edges in these holes expand at a rate of 5 mm / min or more). Can be expanded at a higher percentage relative to diameter), and the edges become thinner, for example as seen in the cross section of the expanded edge of FIG. 91.

Example 20 HER Performance as a Function of Punch Speed and Hole Expansion Speed

Sheets were prepared from alloy 9 according to the atomic ratios provided in Table 2. The slab produced by continuous casting was hot rolled with hot bands and then cold rolled into a sheet about 1.4 mm thick and processed by an annealing cycle. The microstructure of the prepared sheet using both SEM and etched optical microscopy is demonstrated in FIG. 85 showing a typical recrystallized modal structure.

85A and 85B, SEM micrographs show micron-scale properties of austenite-based grains containing some annealing twins and stacking defects. In FIGS. 85C and 85D, the etched samples were examined by light microscopy. It can be seen that the grain boundaries are preferentially etched and the microstructures indicate grain boundaries. Grain size is measured by the line intercept method and ranges from 6 μm to 22 μm with an average value of 15 μm.

Sheets with recrystallized modal structure were used for the HER test. Specimens of 89 × 89 mm size were cut from the sheet with wire EDM. 10 mm diameter holes were punched at two different speeds (5 mm / sec using Instron mechanical test frames and 228 mm / sec using commercial punch presses), flat punch shapes and about 12.5% and 16, respectively Punch spacing for die of% was used.

Hole expansion rate (HER) testing was performed on an Interlaken Technologies SP-225 hydraulic press and consisted of raising a conical punch that evenly extended the hole radially outward. Two hole expansion speeds of 3 mm / min and 50 mm / min, ie conical ram movement speed, were used. The digital image camera system focused on conical punches and monitored the edges of the holes to look for evidence of crack formation and propagation.

Initial diameters of the holes were measured twice using calipers, measurements were made in 90 ° increments and averaged to obtain initial hole diameters. The conical punch was continuously raised until a crack propagated through the specimen thickness. The hole expansion rate was calculated as a percentage of the initial hole diameter measured before the start of the test at the time the test was stopped. Four diameter measurements were made using a caliper every 45 ° after expansion and averaged to account for the asymmetry of the holes due to cracking.

The hole expansion rate values for the test are reported in Table 49. In samples with holes punched at 5 mm / sec, the HER value is in the range of 2.4 to 18.5%. For the 228 mm / second hole punching speed, the HER value is significantly higher in the range of 33.8 to 75.0%. The influence of the expansion speed is illustrated in FIG. 86. Increasing the expansion speed results in higher HER values regardless of the punch speeds used (ie, 5 mm / sec and 228 mm / sec).

TABLE 49 Hole Expansion Rate of Alloy 9 Sheets at Different Punching and Expansion Rates

Punch Spacing (%) Hole punch speed
(mm / sec) Hole expansion speed
(mm / min) HER
(%) 16 228 3 33.8 16 228 3 41.3 16 228 50 63.1 16 228 50 75.0 12.5 5 3 2.4 12.5 5 3 7.9 12.5 5 50 12.7 12.5 5 50 18.5

Fischer Feritscope FMP30 was used to measure the percent magnetic phase volume (Fe%) of the HER tested samples at different hole punching rates and hole expansion rates. The results are reported in Table 50. 87 illustrates the effect of the percent magnetic field volume of the sample tested as a function of distance from the hole edge. As can be seen at higher punch speeds and / or higher expansion speeds, after the test is finished, the magnetic phase volume% increases near the hole edge and also moves away from the hole edge into the material. The magnetic phase volume (Fe%) is consistent with an increase in the amount of structure # 5 in Table 1 formed during deformation due to the formation of magnetic nanoscale α-iron from the starting nonmagnetic austenite present in structure # 4.

TABLE 50 Magnetic phase volume (Fe%) of alloy 9 at different hole punching speeds and hole expansion rates as a function of distance from hole edge after expansion

Hole Creation and Expansion Parameters Hole Punching Speed (mm / sec) 228 228 228 228 5 5 5 5 Punch Spacing (%) 16 16 16 16 12.5 12.5 12.5 12.5 Hole Expansion Speed (mm / min) 3 3 50 50 3 3 50 50 HER (%) 33.8 41.3 63.1 75.0 2.4 7.9 12.7 18.5 Distance from hole (mm) Magnetic phase volume% (Fe%) One 27.3 31.6 37.9 39.1 7.1 9.5 13.5 20.1 2.5 17 21.1 29.6 36 2.4 2.7 6.5 6.2 4 6 7.5 17.4 24.6 0.94 1.1 2.4 2.4 5.5 2.2 2.8 6.3 11.3 0.47 0.45 0.96 0.75 7 0.82 0.89 2.8 4.4 0.21 0.29 0.38 0.28 8.5 0.33 0.35 1.3 1.9 0.23 0.22 0.24 0.16 10 0.21 0.21 0.66 1.1 0.21 0.2 0.2 0.13 11.5 0.15 0.16 0.42 0.67 0.2 0.18 0.21 0.12 13 0.13 0.14 0.26 0.37 0.18 0.18 0.22 0.11 14.5 0.12 0.13 0.25 0.31 0.19 0.18 0.23 0.11 16 0.13 0.14 0.31 0.38 0.19 0.19 0.22 0.13 17.5 0.2 0.22 0.53 0.84 0.19 0.2 0.24 0.14 19 0.16 0.25 0.37 0.61 0.2 0.22 0.22 0.12 22 0.11 0.13 0.21 0.24 0.19 0.21 0.22 0.1 25 0.12 0.12 0.19 0.23 0.19 0.2 0.2 0.11

This embodiment illustrates that the relative resistance to cracking of the edges can be increased by increasing the hole punching speed, hole expansion speed, or both, in an exemplary case of forming an edge in a hole. Sheets from alloy 9 tested in this example may increase the hole expansion rate by increasing the hole punching speed (ie, 5 to 228 mm / sec) and / or the hole expansion speed (ie, 3 to 50 mm / min). Proved. Thus, preferably for the alloy of the subject matter herein, an edge is formed in the alloy and the edge is expanded at a rate of 5 mm / min or more. The percent magnetic phase volume (Fe%) of the tested samples increases with increasing hole punching rate and / or hole expanding rate over the investigated range. Higher local formability and edge in the material as measured by HER, with a relatively higher amount of deformation available at and near the hole edge during the increased hole punching speed or hole expansion speed disclosed herein. Resistance to cracking is achieved.

Example # 21 HER Performance as a Function of the Hole Preparation Method

Slabs with a thickness of 50 mm were lab cast into three selected alloys according to the atomic ratios provided in Table 2, and hot rolled, cold rolled and annealed at 850 ° C. for 10 minutes as described herein. Sheets from each alloy with recrystallized modal structures with a final thickness of 1.2 mm were used to demonstrate the effect of hole expansion rate on HER performance.

Specimens of 89 × 89 mm size were cut from the sheet with wire EDM. 10 mm diameter holes were prepared by various methods including punching, EDM cutting, milling, and laser cutting. Hole punching was performed using a Komatsu OBS80-3 press at a low quasistatic punching speed of 0.25 mm / sec with a punch gap of 16% of the die. EDM cut holes were subjected to final cuts as a parameter to form a visually smooth surface first after rough cuts. During hole milling, the holes were pilot drilled, reamed to size, and then deburred. Laser cut samples were cut on a 4 kW fiber optic Mazak Optiplex 4020 Fiber II machine.

Hole expansion rate (HER) testing was performed on an Interlaken Technologies SP-225 hydraulic press and consisted of raising a conical punch that evenly extended the hole radially outward. In FIG. 88, the results of the HER test are provided as a function of the hole preparation method for each alloy. As shown, for punched holes, the HER value is the lowest in the three alloys in the range of 6 to 12%. Samples with EDM cutting holes, milling holes and laser cutting holes exhibit high HER values of 65-140% +. About 140% expansion represents the maximum extension limit of the press crosshead under test, so that samples with EDM cut holes from alloy 12 and samples with milled holes from alloy 9 and alloy 12 did not reach the expansion limit during the HER test. Note (ie actual value> 140%).

In FIG. 89, SEM images of the sample cross section near the hole edges before expansion are provided at low magnification for samples from Alloy 1 with holes prepared in different ways. In the punched sample (FIG. 89A), one can see the typical rollover area at the top and the burr area at the bottom. In addition, a hemispherical shear influence region can be seen at the edge of the hole and the deepest penetration point is about 0.5 mm. Similar shear influence zones were similarly observed in punched samples of the other two alloys, but not in samples with holes formed by the non-punching method. Note that all methods used for hole preparation introduced some kind of defects at the hole edges. In the EDM cutting hole (FIG. 89B), the edges are vertical in the cross-sectional image but small micron-scale cutting defects can be seen at the surface, and in the milled sample (FIG. 89C), the edges of the holes are trapezoidal and laser cut holes ( In FIG. 89 d), the edges deviated laterally as the laser penetrated the sample. 90 shows hole edges before expansion at higher magnification for samples of alloy 1 with holes prepared by different methods including punching at a hole punching speed of 0.25 mm / sec, EDM cutting holes, milled holes, and laser cutting holes. SEM images of the cross section in the vicinity (at the edge and up to 0.7 mm from the edge) are provided. The microstructures near the hole edges are shown in FIGS. 90A, 90B, 90C and 90D, respectively. As can be seen, the edge of the hole punched at 0.25 mm / sec (FIG. 90A) is relatively highly deformed so that a low HER value is observed. This structure near the edge of the punched sample represents the refined high-strength nanomodal structure of structure # 5 of Table 1, and the structures near the hole edges of the EDM cut holes, milled holes, and laser cut holes are shown in Structure # Represents a recrystallized modal structure of 4. However, in the embodiment where the holes are formed by the non-punching method (Figs. 90B, 90C, 90D), the alloy obtained has a high HER value of 65 to 140% +, which is consistent with the malleability of structure # 4 near the hole edge. Has excellent local formability. In FIGS. 91A (punched holes), FIG. 91B (EDM cut holes), FIG. 92C (milled holes), and FIG. 91D (laser cut holes), SEM images of cross sections near the hole edges after the HER test are samples of alloy 1 It is provided at low magnification for. Since the expansion of the hole makes the sample thinner near the hole edge, the thickness of the sample near the hole is smaller in the expanded sample with the higher HER value.

In FIG. 92, after the HER test (ie after expansion to fracture by crack) at higher magnification for samples of Alloy 1 with holes prepared in different ways showing similarly deformed structures in all cases near the hole edges. An image of the sample cross section is provided. Since the hole expansion and deformation of the edges have been completed, the microstructures near all hole edges are similar, which represents the refined high strength nanomodal structure of structure # 5 in Table 1. This example is a local obtained from the alloys herein. The effect of edge preparation on the phosphor formability is shown. Slow punching of 0.25 mm / sec results in structural changes near the hole edge, which is consistent with previous embodiments leading to limited local formability of the edge and low HER values. However, in the embodiment where the holes are formed by the non-punching method, the alloy obtained has excellent local formability with a high HER value of 65 to 140% +, consistent with the malleable properties at the sample and hole edges.

Claims (20)

As a method of expanding the edge of the alloy,
a. Supplying a metal alloy containing at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, melting the alloy, and at a rate of 250 K / sec or less Cooling or solidifying to a thickness of 2.0 mm to 500 mm and forming an alloy with Tm;
b. The alloy is heated to a temperature of 700 ° C. to less than the Tm of the alloy, the thickness of the alloy is reduced at a strain rate of 10 ⁻⁶ to 10 ⁴ , a maximum tensile strength of 921 MPa to 1413 MPa and 12.0% to 77.7% Providing a resultant first alloy having an elongation of;
c. Stressing the resulting first alloy to provide a resulting second alloy having a maximum tensile strength of 1356 MPa to 1831 MPa and an elongation of 1.6% to 32.8%;
d. Heating the resulting second alloy to a temperature below Tm and forming the resulting third alloy having an elongation of 6.6% to 86.7%;
e. Forming an edge in the resulting alloy, and extending the edge at a speed of at least 5 mm / min. The method of claim 1,
Wherein the edge extends at a speed ranging from 5 mm / min to 100 mm / min. The method of claim 1,
Wherein said alloy comprises Fe and at least 5 elements selected from Si, Mn, B, Cr, Bi, Cu or C. The method of claim 1,
Wherein said alloy comprises Fe and at least six elements selected from Si, Mn, B, Cr, Ni, Cu or C. The method of claim 1,
Wherein the alloy comprises Fe, Si, Mn, B, Cr, Ni, Cu, and C. 18. The method of claim 1,
And the yield strength of the alloy of 197 to 1372 MPa is obtained by heating in step (d). The method of claim 1,
Heating at the step (d) results in a maximum tensile strength of the alloy of 799 to 1683 MPa being obtained. The method of claim 1,
Prior to step (e), the edges in the alloy are exposed to temperatures in the temperature range of 400 ° C. to less than the Tm of the alloy. The method of claim 1,
In said step (e), said edge forms an inner hole and / or an outer edge. The method of claim 1,
In step (e), the edge is formed by punching, piercing, drilling, cutting, cropping, EDM cutting, waterjet cutting, laser cutting, or milling. The method of claim 1,
Wherein the edge in the alloy is formed by a progressive die stamping operation. The method of claim 1,
And the extended edge in the alloy is located in the vehicle. The method of claim 1,
Wherein the expanded edge in the alloy is part of a vehicle frame, vehicle chassis, or vehicle panel. As a method of expanding the edge of the alloy,
Supplying a metal alloy comprising at least 50 atomic% iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu or C, the alloy having a maximum tensile strength of 799 MPa to 1683 MPa and Having an elongation of 6.6 to 86.7%;
Forming an edge in the alloy; And
Expanding the edge of the alloy at a rate of at least 5 mm / min. The method of claim 14,
Wherein the edge extends at a rate of 5 mm / min to 100 mm / min. The method of claim 14,
Forming an edge in the alloy is performed by punching at a speed of at least 5 mm / sec. The method of claim 16,
Said punching speed is from 5 mm / sec to 228 mm / sec. The method of claim 14,
The edge is formed in the alloy at a punch speed of 5 mm / sec to 228 mm / sec, and the edge extends at a speed of 5 mm / min to 100 mm / min. The method of claim 14,
And the alloy with the extended edge is located in the vehicle. The method of claim 14,
Wherein the alloy having an extended edge in the alloy is part of a vehicle frame, vehicle chassis, or vehicle panel.

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