Note: Descriptions are shown in the official language in which they were submitted.
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Improved Edge Formability In Metallic Alloys
Cross-Reference to Related Applications
This application claims priority to U.S. Patent Application 15/438,313, filed
February
21, 2017, which is a continuation-in-part of U.S. Patent Application
15/094,554 filed April 8,
2016, which claims the benefit of U.S. Provisional Patent Application Serial
No. 62/146,048
filed on April 10, 2015 and U.S. Provisional Patent Application Serial No.
62/257,070 filed
on November 18, 2015, which is fully incorporated herein by reference.
Field of Invention
This disclosure relates to methods for mechanical property improvement in a
metallic alloy
that has undergone one or more mechanical property losses as a consequence of
shearing,
such as in the formation of a sheared edge portion or a punched hole. More
specifically,
methods are disclosed that provide the ability to improve mechanical
properties of metallic
alloys that have been formed with one or more sheared edges which may
otherwise serve as a
limiting factor for industrial applications.
Background
From ancient tools to modern skyscrapers and automobiles, steel has driven
human
innovation for hundreds of years. Abundant in the Earth's crust, iron and its
associated alloys
have provided humanity with solutions to many daunting developmental barriers.
From
humble beginnings, steel development has progressed considerably within the
past two
centuries, with new varieties of steel becoming available every few years.
These steel alloys
can be broken up into three classes based upon measured properties, in
particular maximum
tensile strain and tensile stress prior to failure. These three classes are:
Low Strength Steels
(LSS), High Strength Steels (HSS), and Advanced High Strength Steels (AHSS).
Low
Strength Steels (LSS) are generally classified as exhibiting ultimate tensile
strengths less than
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270 MPa and include such types as interstitial free and mild steels. High-
Strength Steels
(HSS) are classified as exhibiting ultimate tensile strengths from 270 to 700
MPa and include
such types as high strength low alloy, high strength interstitial free and
bake hardenable
steels. Advanced High-Strength Steels (AHSS) steels are classified by ultimate
tensile
strengths greater than 700 MPa and include such types as Martensitic steels
(MS), Dual Phase
(DP) steels, Transformation Induced Plasticity (TRIP) steels, and Complex
Phase (CP) steels.
As the strength level increases the trend in maximum tensile elongation
(ductility) of the steel
is negative, with decreasing elongation at high ultimate tensile strengths.
For example,
tensile elongation of LSS, HSS and AHSS ranges from 25% to 55%, 10% to 45%,
and 4% to
30%, respectively.
Production of steel continues to increase, with a current US production around
100 million
tons per year with an estimated value of S75 billion. Steel utilization in
vehicles is also high,
with advanced high strength steels (AHSS) currently at 17% and forecast to
grow by 300% in
the coming years [American Iron and Steel Institute. (2013). Profile 2013.
Washington,
D.C.]. With current market trends and governmental regulations pushing towards
higher
efficiency in vehicles, AHSS are increasingly being pursued for their ability
to provide high
strength to mass ratio. The high strength of AHSS allows for a designer to
reduce the
thickness of a finished part while still maintaining comparable or improved
mechanical
properties. In reducing the thickness of a part, less mass is needed to attain
the same or better
mechanical properties for the vehicle thereby improving vehicle fuel
efficiency. This allows
the designer to improve the fuel efficiency of a vehicle while not
compromising on safety.
One key attribute for next generation steels is formability. Formability is
the ability of a
material to be made into a particular geometry without cracking, rupturing or
otherwise
undergoing failure. High formability steel provides benefit to a part designer
by allowing for
the creation of more complex part geometries allowing for reduction in weight.
Formability
may be further broken into two distinct forms: edge formability and bulk
formability. Edge
formability is the ability for an edge to be formed into a certain shape.
Edges on materials
are created through a variety of methods in industrial processes, including
but not limited to
punching, shearing, piercing, stamping, perforating, cutting, or cropping.
Furthermore, the
devices used to create these edges are as diverse as the methods, including
but not limited to
various types of mechanical presses, hydraulic presses, and/or electromagnetic
presses.
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Depending upon the application and material undergoing the operation, the
range of speeds
for edge creation is also widely varying, with speeds as low as 0.25 mm/s and
as high as 3700
mm/s. The wide variety of edge forming methods, devices, and speeds results in
a myriad of
different edge conditions in use commercially today.
Edges, being free surfaces, are dominated by defects such as cracks or
structural changes in
the sheet resulting from the creation of the sheet edge. These defects
adversely affect the
edge formability during forming operations, leading to a decrease in effective
ductility at the
edge. Bulk formability on the other hand is dominated by the intrinsic
ductility, structure,
and associated stress state of the metal during the forming operation. Bulk
formability is
affected primarily by available deformation mechanisms such as dislocations,
twinning, and
phase transformations. Bulk formability is maximized when these available
deformation
mechanisms are saturated within the material, with improved bulk formability
resulting from
an increased number and availability of these mechanisms.
Edge formability can be measured through hole expansion measurements, whereby
a hole is
made in a sheet and that hole is expanded by means of a conical punch.
Previous studies
have shown that conventional AHSS materials suffer from reduced edge
formability
compared with other LSS and HSS when measured by hole expansion [M.S. Billur,
T.
Altan, "Challenges in forming advanced high strength steels", Proceedings of
New
Developments in Sheet Metal Forming, pp.285-304, 20121. For example, Dual
Phase (DP)
steels with ultimate tensile strength of 780 MPa achieve less than 20% hole
expansion,
whereas Interstitial Free steels (IF) with ultimate tensile strength of
approximately 400 MPa
achieve around 100% hole expansion ratio. This reduced edge formability
complicates
adoption of AHSS in automotive applications, despite possessing desirable bulk
formability.
Summary
A method for improving one or more mechanical properties in a metallic alloy
that has
undergone a mechanical property loss as a consequence of the formation of one
or more
sheared edges comprising:
a. supplying a metal alloy comprising at least 50 atomic % iron and at least
four
or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said
alloy
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and cooling at a rate of < 250 K/s or solidifying to a thickness of > 2.0 mm
up to
500 mm and forming an alloy having a Tm and matrix grains of 2 tm to 10,000
1.1.m;
b. heating said alloy to a temperature of 700 C and below the Tm of said
alloy
and at a strain rate of 10-6 to 104 and reducing said thickness of said alloy
and
providing a first resulting alloy having an ultimate tensile strength of 921
MPa to
1413 MPa;
c. stressing said first resulting alloy and providing a second resulting alloy
having an ultimate tensile strength of 1356 MPa to 1831 MPa and an elongation
of
1.6% to 32.8%;
d. heating said second resulting alloy to a temperature below Tm and forming a
third resulting alloy having matrix grains of 0.5 tm to 50 tm and having an
elongation (Ei);
e. shearing said alloy and forming one or more sheared edges wherein said
alloy's elongation is reduced to a value of E2 wherein E2 = (0.57 to 0.05)
(E1)
f. reheating said alloy with said one or more sheared edges wherein said
alloy's
reduced elongation observed in step (d) is restored to a level having an
elongation
E3 = (0.48 to 1.21)(E1).
The present disclosure also relates to a method for improving the hole
expansion ratio in a
.. metallic alloy that had undergone a hole expansion ratio loss as a
consequence of forming a
hole with a sheared edge comprising:
a. supplying a metal alloy comprising at least 50 atomic % iron and at least
four
or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said
alloy
and cooling at a rate of < 250 K/s or solidifying to a thickness of > 2.0 mm
up to
500 mm and forming an alloy having a Tm and matrix grains of 2 tm to 10,000
1.1.m;
b. heating said alloy to a temperature of 700 C and below the Tm of said
alloy
and at a strain rate of 10-6 to 104 and reducing said thickness of said alloy
and
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providing a first resulting alloy having an ultimate tensile strength of 921
MPa to
1413 MPa and an elongation of 12.0% to 77.7%;
c. stressing said first resulting alloy and providing a second resulting alloy
having an ultimate tensile strength of 1356 MPa to 1831 MPa and an elongation
of
5 1.6% to 32.8%;
d. heating said second resulting alloy to a temperature of at least 650 C and
below Tm and forming a third resulting alloy having matrix grains of 0.5 um to
50
um and forming a hole therein with shearing wherein said hole has a sheared
edge
and has a first hole expansion ratio (HERD;
e. heating said alloy with said hole and associated HERi wherein said alloy
indicates a second hole expansion ratio (HER2) wherein HER2> HERi.
The present invention also relates to method for improving the hole expansion
ratio in a
metallic alloy that had undergone a hole expansion ratio loss as a consequence
of forming a
hole with a sheared edge comprising:
a. supplying a metal alloy comprising at least 50 atomic % iron and at least
four
or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said
alloy
and cooling at a rate of < 250 K/s or solidifying to a thickness of > 2.0 mm
up to
500 mm and forming an alloy having a Tm and matrix grains of 2 um to 10,000
um;
b. heating said alloy to a temperature of 700 C and below the Tm of said
alloy
and at a strain rate of 10-6 to 104 and reducing said thickness of said alloy
and
providing a first resulting alloy having an ultimate tensile strength of 921
MPa to
1413 MPa and an elongation of 12.0% to 77.7%;
c. stressing said first resulting alloy and providing a second resulting alloy
having an ultimate tensile strength of 1356 MPa to 1831 MPa and an elongation
of
1.6% to 32.8%;
d. heating said second resulting alloy to a temperature of at least 650 C and
below Tm and forming a third resulting alloy having matrix grains of 0.5 um to
50
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um wherein said alloy is characterized as having a first hole expansion ratio
(HERD of 30 to-130% for a hole formed therein without shearing;
e. forming a hole in said second resulting alloy wherein said hole is formed
with
shearing and indicates a second hole expansion ratio (HER2) wherein HER2 =
(0.01 to 0.30)(HER1);
f. heating said alloy wherein HER2 recovers to a value HER3 = (0.60 to 1.0)
HERi.
The present invention also relates to a method for punching one or more holes
in a metallic
alloy comprising:
a. supplying a metal alloy comprising at least 50 atomic % iron and at least
four
or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said
alloy
and cooling at a rate of < 250 K/s or solidifying to a thickness of > 2.0 mm
up to
500 mm and forming an alloy having a Tm and matrix grains of 2 um to 10,000
um;
b. heating said alloy to a temperature of 700 C and below the Tm of said
alloy
and at a strain rate of 10-6 to 104 and reducing said thickness of said alloy
and
providing a first resulting alloy having an ultimate tensile strength of 921
MPa to
1413 MPa and an elongation of 12.0% to 77.7%;
c. stressing said first resulting alloy and providing a second resulting alloy
having an ultimate tensile strength of 1356 MPa to 1831 MPa and an elongation
of
1.6% to 32.8%;
d. heating said second resulting alloy to a temperature of at least 650 C and
below Tm and forming a third resulting alloy having matrix grains of 0.5 um to
50
um and having an elongation (Ei);
e. punching a hole in said alloy at a punch speed of greater than or equal to
10
mm/second wherein said punched hole indicates a hole expansion ratio of
greater
than or equal to 10%.
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The present invention also relates to a method for expanding an edge in an
alloy
a. supplying a metal alloy comprising at least 50 atomic % iron and at least
four
or more elements selected from Si, Mn, B, Cr, Ni, Cu or C and melting said
alloy
and cooling at a rate of < 250 K/s or solidifying to a thickness of > 2.0 mm
up to
500 mm and forming an alloy having a Tm;
b. heating said alloy to a temperature of 700 C and below the Tm of said
alloy
and at a strain rate of 10-6 to 104 and reducing said thickness of said alloy
and
providing a first resulting alloy having an ultimate tensile strength of 921
MPa to
1413 MPa and an elongation of 12.0% to 77.7%;
c. stressing said first resulting alloy and providing a second resulting alloy
having an ultimate tensile strength of 1356 MPa to 1831 MPa and an elongation
of
1.6% to 32.8%;
d. heating said second resulting alloy to a temperature of below Tm and
forming
a third resulting alloy having an elongation of 6.6 % to 86.7%;
e. forming an edge in said resulting alloy and expanding said edge at a speed
of
greater than or equal to 5 mm/min.
The present invention also relates to a method for expanding the edge of an
alloy comprising:
supplying a metal alloy comprising at least 50 atomic % iron and at least four
or
more elements selected from Si, Mn, B, Cr, Ni, Cu or C, wherein said alloy has
an
ultimate tensile strength of 799 MPa to 1683 MPa and an elongation of 6.6 to
86.7%;
forming an edge in said alloy;
expanding said edge in said alloy at a speed of greater than or equal to 5
mm/min.
Brief Description of the Drawings
The detailed description below may be better understood with reference to the
accompanying
FIG.s which are provided for illustrative purposes and are not to be
considered as limiting
any aspect of this invention.
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FIG. 1A Structural pathway for the formation of High Strength Nanomodal
Structure and
associated mechanisms.
FIG. 1B Structural pathway for the formation of Recrystallized Modal Structure
and
Refined High Strength Nanomodal Structure and associated mechanisms.
FIG. 2 Structural pathway toward developing Refined High Strength
Nanomodal
Structure which is tied to industrial processing steps.
FIG. 3 Images of laboratory cast 50 mm slabs from: a) Alloy 9 and b)
Alloy 12.
FIG. 4 Images of hot rolled sheet after laboratory casting from: a)
Alloy 9 and b) Alloy
12.
FIG. 5 Images of cold rolled sheet after laboratory casting and hot
rolling from: a) Alloy
9 and b) Alloy 12.
FIG. 6 Microstructure of solidified Alloy 1 cast at 50 mm thickness: a)
Backscattered
SEM micrograph showing the dendritic nature of the Modal Structure in the as-
cast state, b) Bright-field TEM micrograph showing the details in the matrix
grains, c) Bright-field TEM with selected electron diffraction exhibiting the
ferrite phase in the Modal Structure.
FIG. 7 X-ray diffraction pattern for the Modal Structure in Alloy 1
alloy after
solidification: a) Experimental data, b) Rietveld refinement analysis.
FIG. 8 Microstructure of Alloy 1 after hot rolling to 1.7 mm thickness: a)
Backscattered
SEM micrograph showing the homogenized and refined Nanomodal Structure, b)
Bright-field TEM micrograph showing the details in the matrix grains.
FIG. 9 X-ray diffraction pattern for the Nanomodal Structure in Alloy 1
after hot rolling:
a) Experimental data, b) Rietveld refinement analysis.
FIG. 10 Microstructure of Alloy 1 after cold rolling to 1.2 mm thickness: a)
Backscattered
SEM micrograph showing the High Strength Nanomodal Structure after cold
rolling, b) Bright-field TEM micrograph showing the details in the matrix
grains.
FIG. 11 X-ray diffraction pattern for the High Strength Nanomodal Structure in
Alloy 1
after cold rolling: a) Experimental data, b) Rietveld refinement analysis.
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FIG. 12 Bright-field TEM micrographs of microstructure in Alloy 1 after hot
rolling, cold
rolling and annealing at 850 C for 5 mm exhibiting the Recrystallized Modal
Structure: a) Low magnification image, b) High magnification image with
selected electron diffraction pattern showing crystal structure of austenite
phase.
.. FIG. 13 Backscattered SEM micrographs of microstructure in Alloy 1 after
hot rolling,
cold rolling and annealing at 850 C for 5 min exhibiting the Recrystallized
Modal
Structure: a) Low magnification image, b) High magnification image.
FIG. 14 X-ray diffraction pattern for the Recrystallized Modal Structure in
Alloy 1 after
annealing: a) Experimental data, b) Rietveld refinement analysis.
FIG. 15 Bright-field TEM micrographs of microstructure in Alloy 1 showing
Refined
High Strength Nanomodal Structure (Mixed Microconstituent Structure) formed
after tensile deformation: a) Large grains of untransformed structure and
transformed "pockets" with refined grains; b) Refined structure within a
"pocket".
FIG. 16 Backscattered SEM micrographs of microstructure in Alloy 1 showing
Refined
High Strength Nanomodal Structure (Mixed Microconstituent Structure): a) Low
magnification image, b) High magnification image.
FIG. 17 X-ray diffraction pattern for Refined High Strength Nanomodal
Structure in
Alloy 1 after cold deformation: a) Experimental data, b) Rietveld refinement
analysis.
FIG. 18 Microstructure of solidified Alloy 2 cast at 50 mm thickness: a)
Backscattered
SEM micrograph showing the dendritic nature of the Modal Structure in the as-
cast state, b) Bright-field TEM micrograph showing the details in the matrix
grains.
.. FIG. 19 X-ray diffraction pattern for the Modal Structure in Alloy 2 after
solidification: a)
Experimental data, b) Rietveld refinement analysis.
FIG. 20 Microstructure of Alloy 2 after hot rolling to 1.7 mm thickness: a)
Backscattered
SEM micrograph showing the homogenized and refined Nanomodal Structure, b)
Bright-field TEM micrograph showing the details in the matrix grains.
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FIG. 21 X-ray diffraction pattern for the Nanomodal Structure in Alloy 2
after hot rolling:
a) Experimental data, b) Rietveld refinement analysis.
FIG. 22 Microstructure of Alloy 2 after cold rolling to 1.2 mm thickness: a)
Backscattered
SEM micrograph showing the High Strength Nanomodal Structure after cold
5 rolling, b) Bright-field TEM micrograph showing the details in the
matrix grains.
FIG. 23 X-ray diffraction pattern for the High Strength Nanomodal Structure in
Alloy 2
after cold rolling: a) Experimental data, b) Rietveld refinement analysis.
FIG. 24 Bright-field TEM micrographs of microstructure in Alloy 2 after hot
rolling, cold
rolling and annealing at 850 C for 10 min exhibiting the Recrystallized Modal
10 Structure: a) Low magnification image, b) High magnification image
with
selected electron diffraction pattern showing crystal structure of austenite
phase.
FIG. 25 Backscattered SEM micrographs of microstructure in Alloy 2 after hot
rolling,
cold rolling and annealing at 850 C for 10 mm exhibiting the Recrystallized
Modal Structure: a) Low magnification image, b) High magnification image.
FIG. 26 X-ray diffraction pattern for the Recrystallized Modal Structure in
Alloy 2 after
annealing: a) Experimental data, b) Rietveld refinement analysis.
FIG. 27 Microstructure in Alloy 2 showing Refined High Strength Nanomodal
Structure
(Mixed Microconstituent Structure) formed after tensile deformation: a) Bright-
field TEM micrographs of transformed "pockets" with refined grains; b) Back-
scattered SEM micrograph of the microstructure.
FIG. 28 X-ray diffraction pattern for Refined High Strength Nanomodal
Structure in
Alloy 2 after cold deformation: a) Experimental data, b) Rietveld refinement
analysis.
FIG. 29 Tensile properties of Alloy 1 at various stages of laboratory
processing.
FIG. 30 Tensile results for Alloy 13 at various stages of laboratory
processing.
FIG. 31 Tensile results for Alloy 17 at various stages of laboratory
processing.
FIG. 32 Tensile properties of the sheet in hot rolled state and after
each step of cold
rolling/annealing cycles demonstrating full property reversibility at each
cycle in:
a) Alloy, b) Alloy 2.
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FIG. 33 A bend test schematic showing a bending device with two supports and a
former
(International Organization for Standardization, 2005).
FIG. 34 Images of bend testing samples from Alloy 1 tested to 180 : a) Picture
of a full
set of samples tested to 180 without cracking, and b) A close-up view of the
bend of a tested sample.
FIG. 35 a) Tensile test results of the punched and EDM cut specimens from
selected
alloys demonstrating property decrease due to punched edge damage, b) Tensile
curves of the selected alloys for EDM cut specimens.
FIG. 36 SEM images of the specimen edges in Alloy 1 after a) EDM cutting and
b)
Punching.
FIG. 37 SEM images of the microstructure near the edge in Alloy 1: a) EDM cut
specimens and b) Punched specimens.
FIG. 38 Tensile test results for punched specimens from Alloy 1 before and
after
annealing demonstrating full property recovery from edge damage by annealing.
Data for EDM cut specimens for the same alloy are shown for reference.
FIG. 39 Example tensile stress-strain curves for punched specimens from Alloy
1 with
and without annealing.
FIG. 40 Tensile stress-strain curves illustrating the response of cold
rolled Alloy 1 to
recovery temperatures in the range between 400 C and 850 C; a) Tensile curves,
b) Yield strength.
FIG. 41 Bright-field TEM images of cold rolled ALLOY 1 samples exhibiting the
highly
deformed and textured High Strength Nanomodal Structure: a) Lower
magnification image, b) Higher magnification image.
FIG. 42 Bright-field TEM images of ALLOY 1 samples annealed at 450 C 10 min
exhibiting the highly deformed and textured High Strength Nanomodal Structure
with no recrystallization occurred: a) Lower magnification image, b) Higher
magnification image.
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FIG. 43 Bright-field TEM images of ALLOY 1 samples annealed at 600 C 10 min
exhibiting nanoscale grains signaling the beginning of recrystallization: a)
Lower
magnification image, b) Higher magnification image.
FIG. 44 Bright-field TEM images of ALLOY 1 samples annealed at 650 C 10 min
exhibiting larger grains indicating the higher extent of recrystallization: a)
Lower
magnification image, b) Higher magnification image.
FIG. 45 Bright-field TEM images of ALLOY 1 samples annealed at 700 C 10 min
exhibiting recrystallized grains with a small fraction of untransformed area,
and
electron diffraction shows the recrystallized grains are austenite: a) Lower
magnification image, b) Higher magnification image.
FIG. 46 Model Time Temperature Transformation Diagram representing response of
the
steel alloys herein to temperature at annealing. In the heating curve labeled
A,
recovery mechanisms are activated. In the heating curve labeled B, both
recovery
and recrystallization mechanisms are activated.
FIG. 47 Tensile properties of punched specimens before and after annealing at
different
temperatures: a) Alloy 1, b) Alloy 9, and c) Alloy 12.
FIG. 48 Schematic illustration of the sample position for structural
analysis.
FIG. 49 Alloy 1 punched E8 samples in the as-punched condition: a) Low
magnification
image showing a triangular deformation zone at the punched edge which is
located on the right side of the picture. Additionally close up areas for the
subsequent micrographs are provided, b) Higher magnification image showing
the deformation zone, c) Higher magnification image showing the recrystallized
structure far away from the deformation zone, d) Higher magnification image
showing the deformed structure in the deformation zone.
FIG. 50 Alloy 1 punched E8 samples after annealing at 650 C for 10 mm: a) Low
magnification image showing the deformation zone at edge, punching in upright
direction. Additionally, close up areas for the subsequent micrographs are
provided: b) Higher magnification image showing the deformation zone, c)
Higher magnification image showing the recrystallized structure far away from
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the deformation zone, d) Higher magnification image showing the recovered
structure in the deformation zone.
FIG. 51 Alloy 1 punched E8 samples after annealing at 700 C for 10 mm: a) Low
magnification image showing the deformation zone at edge, punching in upright
direction. Additionally, close up areas for the subsequent micrographs are
provided, b) Higher magnification image showing the deformation zone, c)
Higher magnification image showing the recrystallized structure far away from
the deformation zone, d) Higher magnification image showing the recrystallized
structure in the deformation zone.
FIG. 52 Tensile properties for specimens punched at varied speeds from: a)
Alloy 1, b)
Alloy 9, c) Alloy 12.
FIG. 53 HER results for Alloy 1 in a case of punched vs milled hole.
FIG. 54 Cutting plan for SEM microscopy and microhardness measurement samples
from
HER tested specimens.
FIG. 55 A schematic illustration of microhardness measurement locations.
FIG. 56 Microhardness measurement profile in Alloy 1 HER tested samples with:
a)
EDM cut and b) Punched holes.
FIG. 57 Microhardness profiles for Alloy 1 in various stages of processing and
forming,
demonstrating the progression of edge structure transformation during hole
punching and expansion.
FIG. 58 Microhardness data for HER tested samples from Alloy 1 with punched
and
milled holes. Circles indicate a position of the TEM samples in respect to
hole
edge.
FIG. 59 Bright field TEM image of the microstructure in the Alloy 1 sheet
sample before
HER testing.
FIG. 60 Bright field TEM micrographs of microstructure in the HER test sample
from
Alloy 1 with punched hole (HER = 5%) at a location of - 1.5 mm from the hole
edge: a) main untransformed structure; b) "pocket" of partially transformed
structure.
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FIG. 61 Bright field TEM micrographs of microstructure in the HER test sample
from
Alloy 1 with milled hole (HER = 73.6%) at a location of ¨1.5 mm from the hole
edge in different areas: a) & b).
FIG. 62 Focused Ion Beam (FIB) technique used for precise sampling near the
edge of the
punched hole in the Alloy 1 sample: a) FIB technique showing the general
sample location of the milled TEM sample, b) Close up view of the cut-out TEM
sample with indicated location from the hole edge.
FIG. 63 Bright field TEM micrographs of microstructure in the sample from
Alloy 1 with
a punched hole at a location of ¨10 micron from the hole edge.
FIG. 64 Hole expansion ratio measurements for Alloy 1 with and without
annealing of
punched holes.
FIG. 65 Hole expansion ratio measurements for Alloy 9 with and without
annealing of
punched holes.
FIG. 66 Hole expansion ratio measurements for Alloy 12 with and without
annealing of
punched holes.
FIG. 67 Hole expansion ratio measurements for Alloy 13 with and without
annealing of
punched holes.
FIG. 68 Hole expansion ratio measurements for Alloy 17 with and without
annealing of
punched holes.
FIG. 69 Tensile performance of Alloy 1 tested with different edge
conditions. Note that
tensile samples with Punched edge condition have reduced tensile performance
when compared to tensile samples with wire EDM cut and punched with
subsequent annealing (850 C for 10 minutes) edge conditions.
FIG. 70 Edge formability as measured by hole expansion ratio response of Alloy
1 as a
function of edge condition. Note that holes in the Punched condition have
lower
edge formability than holes in the wire EDM cut and punched with subsequent
annealing (850 C for 10 minutes) conditions.
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FIG. 71 Punch speed dependence of Alloy 1 edge formability as a function
of punch
speed, measured by hole expansion ratio. Note the consistent increase in hole
expansion ratio with increasing punch speed.
FIG. 72 Punch speed dependence of Alloy 9 edge formability as a function of
punch
5 speed, measured by hole expansion ratio. Note the rapid increase in
hole
expansion ratio up to approximately 25 mm/s punch speed followed by a gradual
increase in hole expansion ratio.
FIG. 73 Punch speed dependence of Alloy 12 edge formability as a function of
punch
speed, measured by hole expansion ratio. Note the rapid increase in hole
10 expansion ratio up to approximately 25 mm/s punch speed followed by a
continued increase in hole expansion ratio with punch speeds of >100 mm/s.
FIG. 74 Punch speed dependence of commercial Dual Phase 980 steel edge
formability
measured by hole expansion ratio. Note the hole expansion ratio is
consistently
21% with 3% variance for commercial Dual Phase 980 steel at all punch speeds
15 tested.
FIG. 75 Schematic drawings of non-flat punch geometries: 6 taper (left),
7 conical
(center), and conical flat (right). All dimensions are in millimeters.
FIG. 76 Punch geometry effect on Alloy 1 at 28 mm/s, 114 mm/s, and 228 mm/s
punch
speed. Note that for the Alloy 1, the effect of punch geometry diminishes at
228
mm/s punch speed.
FIG. 77 Punch geometry effect on Alloy 9 at 28 mm/s, 114 mm/s, and 228 mm/s
punch
speeds. Note that the 7 conical punch and the conical flat punch result in
the
highest hole expansion ratio.
FIG. 78 Punch geometry effect on Alloy 12 at 28 mm/s, 114 mm/s, and 228 mm/s
punch
speed. Note that the 7 conical punch results at 228 mm/s punch speed in the
highest hole expansion ratio measured for all alloys.
FIG. 79 Punch geometry effect on Alloy 1 at 228 mm/s punch speed. Note that
all punch
geometries result in nearly equal hole expansion ratios of approximately 21%.
FIG. 80 Hole punch speed dependence of commercial steel grades edge
formability
measured by hole expansion ratio.
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FIG. 81 The
post uniform elongation and hole expansion ratio correlation as predicted by
[Paul S.K., J Mater Eng Perform 2014; 23:3610.1 with data for selected
commercial steel grades from the same paper along with Alloy 1 and Alloy 9
data.
FIG. 82 The measured hole expansion ratio in samples from Alloy 1 as a
function of hole
expansion speed.
FIG. 83 The
measured hole expansion ratio in samples from Alloy 9 as a function of hole
expansion speed.
FIG. 84 The measured hole expansion ratio in samples from Alloy 12 as a
function of
hole expansion speed.
FIG. 85
Images of the microstructure in the sheet from Alloy 9; a) SEM image of the
microstructure, b) Higher magnification SEM image of the microstructure, c)
Optical image of the etched surface, and d) Higher magnification optical image
of the etched surface.
FIG. 86 The measured hole expansion ratio as a function of hole punching speed
and hole
expansion speed for sheet of Alloy 9.
FIG. 87 The average magnetic phases volume percent (Fe%) in the HER tested
samples
with different hole punching speed and hole expansion speed as a function of
the
distance from the hole edge.
FIG. 88 The measured hole expansion ratio in samples from Alloy 1, Alloy 9,
and Alloy
12 as a function of hole preparation method.
FIG. 89 SEM
images at low magnification of the cross section near the hole edge in the
Alloy 1 samples with holes prepared by different methods prior to expansion;
a)
Punched hole, b) EDM cut hole, c) Milled hole, and d) Laser cut hole.
FIG. 90 SEM images at high magnification of the cross section near the hole
edge in the
Alloy 1 samples with holes prepared by different methods prior to expanding at
high magnification; a) Punched hole, b) EDM cut hole, c) Milled hole, and d)
Laser cut hole.
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FIG. 91 SEM
images at low magnification of the cross section near the hole edge in the
Alloy 1 samples with holes prepared by different methods after expansion
during
HER testing; a) Punched hole, b) EDM cut hole, c) Milled hole, and d) Laser
cut
hole.
FIG. 92 SEM images of sample cross sections near the hole edge after HER
testing (i.e.
after expansion until failure by cracking) are provided at higher
magnification for
samples from Alloy 1 with holes prepared by different methods; a) Punched
hole,
b) EDM cut hole, c) Milled hole, and d) Laser cut hole.
Detailed Description
Structures And Mechanisms
The steel alloys herein undergo a unique pathway of structural formation
through specific
mechanisms as illustrated in FIG. 1A and FIG. 1B. Initial structure formation
begins with
melting the alloy and cooling and solidifying and forming an alloy with Modal
Structure
.. (Structure #1, FIG. 1A). The Modal Structure exhibits a primarily
austenitic matrix
(gamma-Fe) which may contain, depending on the specific alloy chemistry,
ferrite grains
(alpha-Fe), martensite, and precipitates including borides (if boron is
present) and/or carbides
(if carbon is present). The grain size of the Modal Structure will depend on
alloy chemistry
and the solidification conditions. For example, thicker as-cast structures
(e.g. thickness of
.. greater than or equal to 2.0 mm) result in relatively slower cooling rate
(e.g. a cooling rate of
less than or equal to 250 K/s) and relatively larger matrix grain size.
Thickness may
therefore preferably be in the range of 2.0 to 500 mm. The Modal Structure
preferably
exhibits an austenitic matrix (gamma-Fe) with grain size and/or dendrite
length from 2 to
10,000 um and precipitates at a size of 0.01 to 5.0 um in laboratory casting.
Matrix grain
size and precipitate size might be larger, up to a factor of 10 at commercial
production
depending on alloy chemistry, starting casting thickness and specific
processing parameters.
Steel alloys herein with the Modal Structure, depending on starting thickness
size and the
specific alloy chemistry typically exhibits the following tensile properties,
yield strength from
144 to 514 MPa, ultimate tensile strength in a range from 411 to 907 MPa, and
total ductility
from 3.7 to 24.4%.
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Steel alloys herein with the Modal Structure (Structure #1, FIG. 1A) can be
homogenized and
refined through the Nanophase Refinement (Mechanism #1, FIG. 1A) by exposing
the steel
alloy to one or more cycles of heat and stress ultimately leading to formation
of the
Nanomodal Structure (Structure #2, FIG. 1A). More specifically, the Modal
Structure, when
formed at thickness of greater than or equal to 2.0 mm, or formed at a cooling
rate of less
than or equal to 250 K/s, is preferably heated to a temperature of 700 C to a
temperature
below the solidus temperature (Tm) and at strain rates of 10-6 to 104 with a
thickness
reduction. Transformation to Structure #2 occurs in a continuous fashion
through the
intermediate Homogenized Modal Structure (Structure #1a, FIG. 1A) as the steel
alloy
undergoes mechanical deformation during successive application of temperature
and stress
and thickness reduction such as what can be configured to occur during hot
rolling.
The Nanomodal Structure (Structure #2, FIG. 1A) has a primary austenitic
matrix (gamma-
Fe) and, depending on chemistry, may additionally contain ferrite grains
(alpha-Fe) and/or
precipitates such as borides (if boron is present) and/or carbides (if carbon
is present).
Depending on starting grain size, the Nanomodal Structure typically exhibits a
primary
austenitic matrix (gamma-Fe) with grain size of 1.0 to 100 um and/or
precipitates at a size 1.0
to 200 nm in laboratory casting. Matrix grain size and precipitate size might
be larger up to a
factor of 5 at commercial production depending on alloy chemistry, starting
casting thickness
and specific processing parameters. Steel alloys herein with the Nanomodal
Structure
typically exhibit the following tensile properties, yield strength from 264 to
574 MPa,
ultimate tensile strength in a range from 921 to 1413 MPa, and total ductility
from 12.0 to
77.7%. Structure #2 is preferably formed at thickness of 1 mm to 500 mm.
When steel alloys herein with the Nanomodal Structure (Structure #2, FIG. 1A)
are subjected
to stress at ambient / near ambient temperature (e.g. 25 C at +/- 5 C), the
Dynamic
Nanophase Strengthening Mechanism (Mechanism #2, FIG. 1A) is activated leading
to
formation of the High Strength Nanomodal Structure (Structure #3, FIG. 1A).
Preferably,
the stress is at a level above the alloy's respective yield strength in a
range from 250 to 600
MPa depending on alloy chemistry. The High Strength Nanomodal structure
typically
exhibits a ferritic matrix (alpha-Fe) which, depending on alloy chemistry, may
additionally
contain austenite grains (gamma-Fe) and precipitate grains which may include
borides (if
boron is present) and/or carbides (if carbon is present). Note that the
strengthening
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transformation occurs during strain under applied stress that defines
Mechanism #2 as a
dynamic process during which the metastable austenitic phase (gamma-Fe)
transforms into
ferrite (alpha-Fe) with precipitates. Note that depending on the starting
chemistry, a fraction
of the austenite will be stable and will not transform. Typically, as low as 5
volume percent
.. and as high as 95 volume percent of the matrix will transform. The High
Strength
Nanomodal Structure typically exhibits a ferritic matrix (alpha-Fe) with
matrix grain size of
25 nm to 50 um and precipitate grains at a size of 1.0 to 200 nm in laboratory
casting. Matrix
grain size and precipitate size might be larger up to a factor of 2 at
commercial production
depending on alloy chemistry, starting casting thickness and specific
processing parameters.
Steel alloys herein with the High Strength Nanomodal Structure typically
exhibits the
following tensile properties, yield strength from 718 to 1645 MPa, ultimate
tensile strength in
a range from 1356 to 1831 MPa, and total ductility from 1.6 to 32.8%.
Structure #3 is
preferably formed at thickness of 0.2 to 25.0 mm.
The High Strength Nanomodal Structure (Structure #3, FIG. 1A and FIG. 1B) has
a
.. capability to undergo Recrystallization (Mechanism #3, FIG. 1B) when
subjected to heating
below the melting point of the alloy with transformation of ferrite grains
back into austenite
leading to formation of Recrystallized Modal Structure (Structure #4, FIG.
1B). Partial
dissolution of nanoscale precipitates also takes place. Presence of borides
and/or carbides is
possible in the material depending on alloy chemistry. Preferred temperature
ranges for a
complete transformation occur from 650 C up to the Tm of the specific alloy.
When
recrystallized, the Structure #4 contains few dislocations or twins and
stacking faults can be
found in some recrystallized grains. Note that at lower temperatures from 400
to 650 C,
recovery mechanisms may occur. The Recrystallized Modal Structure (Structure
#4, FIG.
1B) typically exhibits a primary austenitic matrix (gamma-Fe) with grain size
of 0.5 to 50 um
and precipitate grains at a size of 1.0 to 200 nm in laboratory casting.
Matrix grain size and
precipitate size might be larger up to a factor of 2 at commercial production
depending on
alloy chemistry, starting casting thickness and specific processing
parameters. Steel alloys
herein with the Recrystallized Modal Structure typically exhibit the following
tensile
properties: yield strength from 197 to 1372 MPa, ultimate tensile strength in
a range from 799
to 1683 MPa, and total ductility from 10.6 to 86.7%.
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Steel alloys herein with the Recrystallized Modal Structure (Structure #4,
FIG. 1B) undergo
Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B) upon stressing
above
yield at ambient / near ambient temperature (e.g. 25 C +/- 5 C) that leads to
formation of the
Refined High Strength Nanomodal Structure (Structure #5, FIG. 1B). Preferably
the stress to
5 initiate Mechanism #4 is at a level above yield strength in a range 197
to 1372 MPa. Similar
to Mechanism #2, Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B)
is a
dynamic process during which the metastable austenitic phase transforms into
ferrite with
precipitate resulting generally in further grain refinement as compared to
Structure #3 for the
same alloy. One characteristic feature of the Refined High Strength Nanomodal
Structure
10 (Structure #5, FIG. 1B) is that significant refinement occurs during
phase transformation in
the randomly distributed "pockets" of microstructure while other areas remain
untransformed. Note that depending on the starting chemistry, a fraction of
the austenite will
be stable and the area containing the stabilized austenite will not transform.
Typically, as low
as 5 volume percent and as high as 95 volume percent of the matrix in the
distributed
15 "pockets" will transform. The presence of borides (if boron is present)
and/or carbides (if
carbon is present) is possible in the material depending on alloy chemistry.
The
untransformed part of the microstructure is represented by austenitic grains
(gamma-Fe) with
a size from 0.5 to 50 um and additionally may contain distributed precipitates
with size of 1
to 200 nm. These highly deformed austenitic grains contain a relatively large
number of
20 dislocations due to existing dislocation processes occurring during
deformation resulting in
high fraction of dislocations (108 to 1010 mm-2). The transformed part of the
microstructure
during deformation is represented by refined ferrite grains (alpha-Fe) with
additional
precipitate through Nanophase Refinement & Strengthening (Mechanism #4, FIG.
1B). The
size of refined grains of ferrite (alpha-Fe) varies from 50 to 2000 nm and
size of precipitates
is in a range from 1 to 200 nm in laboratory casting. Matrix grain size and
precipitate size
might be larger up to a factor of 2 at commercial production depending on
alloy chemistry,
starting casting thickness and specific processing parameters. The size of the
"pockets" of
transformed and highly refined microstructure typically varies from 0.5 to 20
um. The
volume fraction of the transformed vs untransformed areas in the
microstructure can be
varied by changing the alloy chemistry including austenite stability from
typically a 95:5
ratio to 5:95, respectively. Steel alloys herein with the Refined High
Strength Nanomodal
Structure typically exhibit the following tensile properties: yield strength
from 718 to 1645
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MPa, ultimate tensile strength in a range from 1356 to 1831 MPa, and total
ductility from 1.6
to 32.8%.
Steel alloys herein with the Refined High Strength Nanomodal Structure
(Structure #5, FIG.
1B) may then be exposed to elevated temperatures leading back to formation of
a
Recrystallized Modal Structure (Structure #4, FIG. 1B). Typical temperature
ranges for a
complete transformation occur from 650 C up to the Tm of the specific alloy
(as illustrated in
FIG. 1B) while lower temperatures from 400 C to temperatures less than 650 C,
activate
recovery mechanisms and may cause partial recrystallization. Stressing and
heating may be
repeated multiple times to achieve desired product geometry including but not
limited to
relatively thin gauges of the sheet, relatively small diameter of the tube or
rod, complex shape
of final part, etc. with targeted properties. Final thicknesses of the
material may therefore fall
in the range from 0.2 to 25 mm. Note that cubic precipitates may be present in
the steel
alloys herein at all stages with a Fm3m (#225) space group. Additional
nanoscale
precipitates may be formed as a result of deformation through Dynamic
Nanophase
Strengthening Mechanism (Mechanism #2) and/or Nanophase Refinement &
Strengthening
(Mechanism #4) that are represented by a dihexagonal pyramidal class hexagonal
phase with
a P63m, space group (#186) and/or a ditrigonal dipyramidal class with a
hexagonal P6bar2C
space group (#190). The precipitate nature and volume fraction depends on the
alloy
composition and processing history. The size of nanoprecipitates can range
from 1 nm to
tens of nanometers, but in most cases below 20 nm. Volume fraction of
precipitates is
generally less than 20%.
Mechanisms During Sheet Production Through Slab Casting
The structures and enabling mechanisms for the steel alloys herein are
applicable to
commercial production using existing process flows. See FIG. 2. Steel slabs
are commonly
produced by continuous casting with a multitude of subsequent processing
variations to get to
the final product form which is commonly coils of sheet. A detailed structural
evolution in
steel alloys herein from
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casting to final product with respect to each step of slab processing into
sheet product is
illustrated in FIG. 2.
The formation of Modal Structure (Structure #1) in steel alloys herein occurs
during alloy
solidification. The Modal Structure may be preferably formed by heating the
alloys herein at
temperatures in the range of above their melting point and in a range of 1100
C to 2000 C
and cooling below the melting temperature of the alloy, which corresponds to
preferably
cooling in the range of 1x103 to 1x10-3 K/s. The as-cast thickness will be
dependent on the
production method with Thin Slab Casting typically in the range of 20 to 150
mm in
thickness and Thick Slab Casting typically in the range of 150 to 500 mm in
thickness.
Accordingly, as cast thickness may fall in the range of 20 to 500 mm, and at
all values
therein, in 1 mm increments. Accordingly, as cast thickness may be 21 mm, 22
mm, 23 mm,
etc., up to 500 mm.
Hot rolling of solidified slabs from the alloys is the next processing step
with production
either of transfer bars in the case of Thick Slab Casting or coils in the case
of Thin Slab
Casting. During this process, the Modal Structure transforms in a continuous
fashion into a
partial and then fully Homogenized Modal Structure (Structure #1a) through
Nanophase
Refinement (Mechanism #1). Once homogenization and resulting refinement is
completed,
the Nanomodal Structure (Structure #2) forms. The resulting hot band coils
which are a
product of the hot rolling process is typically in the range of 1 to 20 mm in
thickness.
Cold rolling is a widely used method for sheet production that is utilized to
achieve targeted
thickness for particular applications. For AHSS, thinner gauges are usually
targeted in the
range of 0.4 to 2 mm. To achieve the finer gauge thicknesses, cold rolling can
be applied
through multiple passes with or without intermediate annealing between passes.
Typical
reduction per pass is 5 to 70% depending on the material properties and
equipment capability.
The number of passes before the intermediate annealing also depends on
materials properties
and level of strain hardening during cold deformation. For the steel alloys
herein, the cold
rolling will trigger Dynamic Nanophase Strengthening (Mechanism #2) leading to
extensive
strain hardening of the resultant sheet and to the formation of the High
Strength Nanomodal
Structure (Structure #3). The properties of the cold rolled sheet from alloys
herein will
depend on the alloy chemistry and can be controlled by the cold rolling
reduction to yield a
fully cold rolled (i.e. hard) product or can be done to yield a range of
properties (i.e. 1/4, 1/2, 3/4
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hard etc.). Depending on the specific process flow, especially starting
thickness and the
amount of hot rolling gauge reduction, often annealing is needed to recover
the ductility of
the material to allow for additional cold rolling gauge reduction.
Intermediate coils can be
annealed by utilizing conventional methods such as batch annealing or
continuous annealing
lines. The cold deformed High Strength Nanomodal Structure (Structure #3) for
the steel
alloys herein will undergo Recrystallization (Mechanism #3) during annealing
leading to the
formation of the Recrystallized Modal Structure (Structure #4). At this stage,
the
recrystallized coils can be a final product with advanced property combination
depending on
the alloy chemistry and targeted markets. In a case when even thinner gauges
of the sheet are
required, recrystallized coils can be subjected to further cold rolling to
achieve targeted
thickness that can be realized by one or multiple cycles of cold rolling /
annealing.
Additional cold deformation of the sheet from alloys herein with
Recrystallized Modal
Structure (Structure #4) leads to structural transformation into Refined High
Strength
Nanomodal Structure (Structure #5) through Nanophase Refinement and
Strengthening
(Mechanism #4). As a result, fully hard coils with final gauge and Refined
High Strength
Nanomodal Structure (Structure #5) can be formed or, in the case of annealing
as a last step
in the cycle, coils of the sheet with final gauge and Recrystallized Modal
Structure (Structure
#4) can also be produced. When coils of recrystallized sheet from alloys
herein utilized for
finished part production by any type of cold deformation such as cold
stamping,
hydroforming, roll forming etc., Refined High Strength Nanomodal Structure
(Structure #5)
will be present in the final product / parts. The final products may be in
many different forms
including sheet, plate, strips, pipes, and tubes and a myriad of complex parts
made through
various metalworking processes.
Mechanisms for Edge Formability
The cyclic nature of these phase transformations going from Recrystallized
Modal Structure
(Structure #4) to Refined High Strength Nanomodal Structure (Structure #5) and
then back to
Recrystallized Modal Structure (Structure #4) is one of the unique phenomenon
and features
of steel alloys herein. As described earlier, this cyclic feature is
applicable during
commercial manufacturing of the sheet, especially for AHSS where thinner gauge
thicknesses
are required (e.g. thickness in the range of 0.2 to 25 mm). Furthermore, these
reversibility
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mechanisms are applicable for the widespread industrial usage of the steel
alloys herein.
While exhibiting exceptional combinations of bulk sheet formability as is
demonstrated by
the tensile and bend properties in this application for the steel alloys
herein, the unique cycle
feature of the phase transformations is enabling for edge formability, which
can be a
significant limiting factor for other AHSS. Table 1 below provides a summary
of the
structure and performance features through stressing and heating cycles
available through
Nanophase Refinement and Strengthening (Mechanism #4). How these structures
and
mechanisms can be harnessed to produce exceptional combinations of both bulk
sheet and
edge formability will be subsequently described herein.
Table 1 Structures and Performance Through Stressing / Heating Cycles
Structure #4 Structure #5
Recrystallized Modal Refined High Strength Nanomodal
Property /
Structure Structure
Mechanism
Untransformed Transformed
"pockets"
Nanophase
Refinement &
Strengthening
Recrystallization mechanism
Structure occurring at elevated Retained austenitic
occurring through
Formation temperatures in cold grains application of
worked material
mechanical stress in
distributed
microstructural
"pockets"
Stress induced
austenite
Recrystallization of cold Precipitation
Transformations transformation into
deformed iron matrix optional
ferrite and
precipitates
Austenite, Ferrite, optionally
Austenite, optionally
Enabling Phases optionally austenite,
ferrite, precipitates
precipitates precipitates
Matrix Grain
0.5 to 50 um 0.5 to 50 um 50 to 2000 nm
Size
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Precipitate Size 1 to 200 nm 1 to 200 nm I 1 to 200 nm
Actual with properties
achieved based on Actual with properties achieved based on
Tensile
formation of the formation of the structure and
fraction of
Response
structure and fraction of transformation
transformation
Yield Strength 197 to 1372 MPa 718 to 1645 MPa
Ultimate Tensile
799 to 1683 MPa 1356 to 1831 MPa
Strength
Total Elongation 6.6 to 86.7% 1.6 to 32.8%
Main Body
The chemical composition of the alloys herein is shown in Table 2 which
provides the
preferred atomic ratios utilized.
5
Table 2 Alloy Chemical Composition
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
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Alloy Fe Cr Ni Mn Cu B Si C
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 alloys herein are iron based metal alloys,
having greater
than or equal to 50 at.% Fe. More preferably, the alloys herein can be
described as
comprising, consisting essentially of, or consisting of the following elements
at the indicated
atomic percent: Fe (61.30 to 83.14 at. %); Si (0 to 7.02 at.%); Mn (0 to 15.86
at.%); B (0 to
6.09 at.%); Cr (0 to 18.90 at.%); Ni (0 to 8.68 at.%); Cu (0 to 2.00 at.%); C
(0 to 3.72 at.%).
In addition, it can be appreciated that the alloys herein are such that they
comprise Fe and at
least four or more, or five or more, or six or more elements selected from Si,
Mn, B, Cr, Ni,
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Cu or C. Most preferably, the alloys herein are such that they comprise,
consist essentially
of, or consist of Fe at a level of 50 at.% or greater along with Si, Mn, B,
Cr, Ni, Cu and C.
Alloy Laboratory Processing
Laboratory processing of the alloys in Table 2 was done to model each step of
industrial
production but on a much smaller scale. Key steps in this process include the
following:
casting, tunnel furnace heating, hot rolling, cold rolling, and annealing.
Casting
Alloys were weighed out into charges ranging from 3,000 to 3,400 grams using
commercially
available ferroadditive powders with known chemistry and impurity content
according to the
atomic ratios in Table 2. Charges were loaded into a zirconia coated silica
crucibles which
was placed into an Indutherm VTC800V vacuum tilt casting machine. The machine
then
evacuated the casting and melting chambers and backfilled with argon to
atmospheric
pressure several times prior to casting to prevent oxidation of the melt. The
melt was heated
with a 14 kHz RF induction coil until fully molten, approximately 5.25 to 6.5
minutes
depending on the alloy composition and charge mass. After the last solids were
observed to
melt it was allowed to heat for an additional 30 to 45 seconds to provide
superheat and ensure
melt homogeneity. The casting machine then evacuated the melting and casting
chambers,
tilted the crucible and poured the melt into a 50 mm thick, 75 to 80 mm wide,
and 125 mm
deep channel in a water cooled copper die. The melt was allowed to cool under
vacuum for
200 seconds before the chamber was filled with argon to atmospheric pressure.
Example
pictures of laboratory cast slabs from two different alloys are shown in FIG.
3.
Tunnel Furnace Heating
Prior to hot rolling, laboratory slabs were loaded into a Lucifer EHS3GT-B18
furnace to heat.
The furnace set point varies between 1100 C to 1250 C depending on alloy
melting point.
The slabs were allowed to soak for 40 minutes prior to hot rolling to ensure
they reach the
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target temperature. Between hot rolling passes the slabs are returned to the
furnace for 4
minutes to allow the slabs to reheat.
Hot Rolling
Pre-heated slabs were pushed out of the tunnel furnace into a Fenn Model 061 2
high rolling
mill. The 50 mm slabs were preferably hot rolled for 5 to 8 passes though the
mill before
being allowed to air cool. After the initial passes each slab had been reduced
between 80 to
85% to a final thickness of between 7.5 and 10 mm. After cooling each
resultant sheet was
sectioned and the bottom 190 mm was hot rolled for an additional 3 to 4 passes
through the
mill, further reducing the plate between 72 to 84% to a final thickness of
between 1.6 and 2.1
mm. Example pictures of laboratory cast slabs from two different alloys after
hot rolling are
shown in FIG. 4.
Cold Rolling
After hot rolling resultant sheets were media blasted with aluminum oxide to
remove the mill
scale and were then cold rolled on a Fenn Model 061 2 high rolling mill. Cold
rolling takes
multiple passes to reduce the thickness of the sheet to a targeted thickness
of typically 1.2
mm. Hot rolled sheets were fed into the mill at steadily decreasing roll gaps
until the
minimum gap is reached. If the material has not yet hit the gauge target,
additional passes at
the minimum gap were used until 1.2 mm thickness was achieved. A large number
of passes
were applied due to limitations of laboratory mill capability. Example
pictures of cold rolled
sheets from two different alloys are shown in FIG. 5.
Annealing
After cold rolling, tensile specimens were cut from the cold rolled sheet via
wire electrical
discharge machining (EDM). These specimens were then annealed with different
parameters
listed in Table 3. Annealing la, lb, 2b were conducted in a Lucifer 7HT-K12
box furnace.
Annealing 2a and 3 was conducted in a Camco Model G-ATM-12FL furnace.
Specimens
which were air normalized were removed from the furnace at the end of the
cycle and
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allowed to cool to room temperature in air. For the furnace cooled specimens,
at the end of
the annealing the furnace was shut off to allow the sample to cool with the
furnace. Note that
the heat treatments were selected for demonstration but were not intended to
be limiting in
scope. High temperature treatments up to just below the melting points for
each alloy are
possible.
Table 3 Annealing Parameters
Annealing Heating Temperature Dwell Cooling Atmosphere
Preheated
la 850 C 5 min Air Normalized Air + Argon
Furnace
Preheated
lb 850 C 10 min Air Normalized Air + Argon
Furnace
45 C/hr to 500 C
Hydrogen +
2a 20 C/hr 850 C 360 min
then Furnace Cool Argon
45 C/hr to 500 C
2b 20 C/hr 850 C 360 min Air
+ Argon
then Air Normalized
3 20 C/hr 1200 C 120 min Furnace Cool
Hydrogen +
Argon
Alloy Properties
Thermal analysis of the alloys herein was performed on as-solidified cast
slabs using a
Netzsch Pegasus 404 Differential Scanning Calorimeter (DSC). Samples of alloys
were
loaded into alumina crucibles which were then loaded into the DSC. The DSC
then
evacuated the chamber and backfilled with argon to atmospheric pressure. A
constant purge
of argon was then started, and a zirconium getter was installed in the gas
flow path to further
reduce the amount of oxygen in the system. The samples were heated until
completely
molten, cooled until completely solidified, then reheated at 10 C/min through
melting.
Measurements of the solidus, liquidus, and peak temperatures were taken from
the second
melting in order to ensure a representative measurement of the material in an
equilibrium
state. In the alloys listed in Table 2, melting occurs in one or multiple
stages with initial
melting from ¨1111 C depending on alloy chemistry and final melting
temperature up to
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-1476 C (Table 4). Variations in melting behavior reflect complex phase
formation at
solidification of the alloys depending on their chemistry.
Table 4 Differential Thermal Analysis Data for Melting Behavior
Solidus Liquidus
Melting Melting Melting
Alloy Temperature
Temperature Peak #1 Peak #2 Peak #3
( C) ( C) ( C) ( C) ( 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
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Solidus Liquidus Melting Melting Melting
Alloy Temperature Temperature Peak #1 Peak #2 Peak #3
( C) ( C) ( C) ( C) ( C)
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 alloys was measured on 9 mm thick sections of hot rolled
material using
the Archimedes method in a specially constructed balance allowing weighing in
both air and
distilled water. The density of each alloy is tabulated in Table 5 and was
found to be in the
range from 7.57 to 7.89 g/cm3. The accuracy of this technique is 0.01 g/cm3.
Table 5 Density of Alloys
Density Density
Alloy Alloy 3
(g/cm3) (g/cm)
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
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Density Density
Alloy Alloy
(g/cm3) (g/cm3)
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 testing frame
using
Instron's Bluehill control software. All tests were conducted at room
temperature, with the
bottom grip fixed and the top grip set to travel upwards at a rate of 0.012
mm/s. Strain data
was collected using Instron's Advanced Video Extensometer. Tensile properties
of the alloys
listed in Table 2 after annealing with parameters listed in Table 3 are shown
below in Table 6
to Table 10. The ultimate tensile strength values may vary from 799 to 1683
MPa with
tensile elongation from 6.6 to 86.7%. The yield strength is in a range from
197 to 978 MPa.
The mechanical characteristic values in the steel alloys herein will depend on
alloy chemistry
and processing conditions. The variation in heat treatment additionally
illustrates the
property variations possible through processing a particular alloy chemistry.
Table 6 Tensile Data for Selected Alloys after Heat Treatment la
Alloy Yield Ultimate Tensile Tensile
Strength Strength Elongation
(MPa) (MPa) (%)
Alloy 1 443 1212 51.1
458 1231 57.9
422 1200 51.9
Alloy 2 484 1278 48.3
485 1264 45.5
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Alloy Yield Ultimate Tensile Tensile
Strength Strength Elongation
(MPa) (MPa) (%)
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
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Alloy Yield Ultimate Tensile Tensile
Strength Strength Elongation
(MPa) (MPa) (%)
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
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Alloy Yield Ultimate Tensile Tensile
Strength Strength Elongation
(MPa) (MPa) (%)
420 1101 44.1
Table 7. Tensile Data for Selected Alloys after Heat Treatment lb
Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
487 1239 57.5
Alloy 1 466 1269 52.5
488 1260 55.8
438 1232 49.7
Alloy 2 431 1228 49.8
431 1231 49.4
522 1172 62.6
Alloy 9 466 1170 61.9
462 1168 61.3
471 1115 63.3
Alloy 12 458 1102 69.3
454 1118 69.1
452 1408 40.5
Alloy 13 435 1416 42.5
432 1396 46.0
448 1132 64.4
Alloy 14 443 1151 60.7
436 1180 54.3
444 1077 66.9
Alloy 15 438 1072 65.3
423 1075 70.5
433 1084 67.5
Alloy 16 432 1072 66.8
423 1071 67.8
Alloy 17 420 946 74.6
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Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
421 939 77.0
425 961 74.9
496 1124 67.4
Alloy 19 434 1118 64.8
435 1117 67.4
434 1154 58.3
Alloy 20 457 1188 54.9
448 1187 60.5
421 1201 54.3
Alloy 21 427 1185 59.9
431 1191 47.8
554 1151 23.5
Alloy 24 538 1142 24.3
562 1151 24.3
500 1274 16.0
Alloy 25 502 1271 15.8
483 1280 16.3
697 1215 20.6
Alloy 26 723 1187 21.3
719 1197 21.5
538 1385 20.6
Alloy 27 574 1397 20.9
544 1388 21.8
978 1592 6.6
Alloy 33 896 1596 7.2
953 1619 7.5
467 1227 56.7
Alloy 34 476 1232 52.7
462 1217 51.6
439 1166 56.3
Alloy 35 438 1166 59.0
440 1177 58.3
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Table 8 Tensile Data for Selected Alloys after Heat Treatment 2a
Yield Ultimate Tensile Tensile
Alloy
Strength Strength Elongation
(MPa) (MPa) (%)
367 1174 46.2
Alloy 2 369 1193 45.1
367 1179 50.2
391 1118 55.7
Alloy 30 389 1116 60.5
401 1113 59.5
413 878 17.6
Alloy 32 399 925 20.5
384 962 21.0
301 1133 37.4
Alloy 31 281 1125 38.7
287 1122 39.0
Table 9 Tensile Data for Selected Alloys after Heat Treatment 2b
Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
396 1093 31.2
Alloy 1 383 1070 30.4
393 1145 34.7
378 1233 49.4
Alloy 2 381 1227 48.3
366 1242 47.7
388 1371 41.3
Alloy 3
389 1388 42.6
335 1338 21.7
Alloy 4 342 1432 30.1
342 1150 17.3
Alloy 5 568 1593 15.2
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Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
595 1596 13.1
735 1605 14.6
399 1283 17.5
Alloy 6 355 1483 24.8
386 1471 23.8
605 1622 16.3
Alloy 7
639 1586 15.2
595 1585 13.6
Alloy 8 743 1623 14.1
791 1554 13.9
381 1125 53.3
Alloy 9 430 1111 44.8
369 1144 51.1
362 1104 37.8
Alloy 10
369 1156 43.5
397 1103 52.4
Alloy 11 390 1086 50.9
402 1115 50.4
358 1055 64.7
Alloy 12 360 1067 64.4
354 1060 62.9
362 982 17.3
Alloy 13 368 961 16.3
370 989 17.0
385 1165 59.0
Alloy 14 396 1156 55.5
437 1155 57.9
357 1056 70.3
Alloy 15 354 1046 68.2
358 1060 70.7
375 1094 67.6
Alloy 16
384 1080 63.4
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Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
326 1054 65.2
368 960 77.2
Alloy 17 370 955 77.9
358 951 75.9
326 1136 17.3
Alloy 18 338 1192 19.1
327 1202 18.5
386 1134 64.5
Alloy 19 378 1100 60.5
438 1093 52.5
386 1172 56.2
Alloy 20 392 1129 42.0
397 1186 57.8
Alloy 21 363 1141 49.0
335 1191 45.7
Alloy 22 322 1189 41.5
348 1168 34.5
398 1077 44.3
Alloy 23
367 1068 44.8
476 1149 28.0
Alloy 24 482 1154 25.9
495 1145 26.2
452 1299 16.0
Alloy 25 454 1287 15.8
441 1278 15.1
619 1196 26.6
Alloy 26 615 1189 26.2
647 1193 26.1
459 1417 17.3
Alloy 27 461 1410 16.8
457 1410 17.1
Alloy 28 507 879 52.3
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Yield Ultimate Tensile Tensile
Alloy Strength Strength Elongation
(MPa) (MPa) (%)
498 874 42.5
493 880 44.7
256 1035 42.3
Alloy 32 257 1004 42.1
257 1049 34.8
830 1494 8.4
Alloy 33 862 1521 8.1
877 1519 8.8
388 1178 59.8
Alloy 34 384 1197 57.7
370 1177 59.1
367 1167 58.5
Alloy 35 369 1167 58.4
375 1161 59.7
Table 10 Tensile Data for Selected Alloys after Heat Treatment 3
Yield
Ultimate Tensile Tensile
Alloy Strength
Strength (MPa) Elongation (%)
(MPa)
238 1142 47.6
Alloy 1 233 1117 46.3
239 1145 53.0
266 1338 38.5
Alloy 3 N/A 1301 37.7
N/A 1291 35.6
N/A 1353 27.7
Alloy 4 N/A 1337 26.1
N/A 1369 29.0
511 1462 12.5
Alloy 5
558 1399 10.6
Alloy 6 311 1465 24.6
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Yield
Ultimate Tensile Tensile
Alloy Strength
Strength (MPa) Elongation (%)
(MPa)
308 1467 21.8
308 1460 25.0
727 1502 12.5
Alloy 7 639 1474 11.3
685 1520 12.4
700 1384 12.3
Alloy 8
750 1431 13.3
234 1087 55.0
Alloy 9 240 1070 56.4
242 1049 58.3
229 1073 50.6
Alloy 10 228 1082 56.5
229 1077 54.2
232 1038 63.8
Alloy 11 232 1009 62.4
228 999 66.1
229 979 65.6
Alloy 12 228 992 57.5
222 963 66.2
277 1338 37.3
Alloy 13 261 1352 35.9
272 1353 34.9
228 1074 58.5
Alloy 14 239 1077 54.1
230 1068 49.1
206 991 60.9
Alloy 15
208 1024 58.9
199 1006 57.7
Alloy 16 242 987 53.4
208 995 57.0
Alloy 17 222 844 72.6
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Yield
Ultimate Tensile Tensile
Alloy Strength
Strength (MPa) Elongation (%)
(MPa)
197 867 64.9
213 869 66.5
288 1415 32.6
Alloy 18 300 1415 32.1
297 1421 29.6
225 1032 58.5
Alloy 19 213 1019 61.1
214 1017 58.4
233 1111 57.3
Alloy 20 227 1071 53.0
230 1091 49.4
238 1073 50.6
Alloy 21 228 1069 56.5
246 1110 52.0
217 1157 47.0
Alloy 22 236 1154 46.8
218 1154 47.7
208 979 45.4
Alloy 23 204 984 43.4
204 972 38.9
277 811 86.7
Alloy 28 279 802 86.0
277 799 82.0
203 958 33.3
Alloy 32 206 966 39.5
210 979 36.3
216 1109 52.8
Alloy 34 230 1144 55.9
231 1123 52.3
230 1104 51.7
Alloy 35
231 1087 59.0
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Yield
Ultimate Tensile Tensile
Alloy Strength
Strength (MPa) Elongation (%)
(MPa)
220 1084 54.4
Case Examples
Case Example #1: Structural Development Pathway in Alloy 1
A laboratory slab with thickness of 50 mm was cast from Alloy 1 that was then
laboratory
processed by hot rolling, cold rolling and annealing at 850 C for 5 mm as
described in Main
Body section of current application. Microstructure of the alloy was examined
at each step of
processing by SEM, TEM and x-ray analysis.
For SEM study, the cross section of the slab samples was ground on SiC
abrasive papers with
reduced grit size, and then polished progressively with diamond media paste
down to 1 um.
The final polishing was done with 0.02 um grit SiO2 solution. Microstructures
were
examined by SEM using an EVO-MA10 scanning electron microscope manufactured by
Carl
Zeiss SMT Inc. To prepare TEM specimens, the samples were first cut by EDM,
and then
thinned by grinding with pads of reduced grit size every time. Further
thinning to make foils
of 60 to 70 um thickness was done by polishing with 9 um, 3 um and 1 um
diamond
suspension solution respectively. Discs of 3 mm in diameter were punched from
the foils and
the final polishing was completed with electropolishing using a twin-jet
polisher. The
chemical solution used was a 30% nitric acid mixed in methanol base. In case
of insufficient
thin area for TEM observation, the TEM specimens may be ion-milled using a
Gatan
Precision Ion Polishing System (PIPS). The ion-milling usually is done at 4.5
keV, and the
inclination angle is reduced from 4 to 2 to open up the thin area. The TEM
studies were
done using a JEOL 2100 high-resolution microscope operated at 200 kV. X-ray
diffraction
was done using a PANalytical X'Pert MPD diffractometer with a Cu Koc x-ray
tube and
operated at 45 kV with a filament current of 40 mA. Scans were run with a step
size of 0.01
and from 250 to 950 two-theta with silicon incorporated to adjust for
instrument zero angle
.. shift. The resulting scans were then subsequently analyzed using Rietveld
analysis using
Siroquant software.
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Modal Structure was formed in the Alloy 1 slab with 50 mm thickness after
solidification.
The Modal Structure (Structure #1) is represented by a dendritic structure
that is composed of
several phases. In FIG. 6a, the backscattered SEM image shows the dendritic
arms that are
shown in dark contrast while the matrix phase is in bright contrast. Note that
small casting
pores are found as exhibited (black holes) in the SEM micrograph. TEM studies
show that
the matrix phase is primarily austenite (gamma-Fe) with stacking faults (FIG.
6b). The
presence of stacking faults indicates a face-centered-cubic structure
(austenite). TEM also
suggests that other phases could be formed in the Modal Structure. As shown in
FIG. 6c, a
dark phase is found that identified as a ferrite phase with body-centered
cubic structure
(alpha-Fe) according to selected electron diffraction pattern. X-ray
diffraction analysis shows
that the Modal Structure of the Alloy 1 contains austenite, ferrite, iron
manganese compound
and some martensite (FIG. 7). Generally, austenite is the dominant phase in
the Alloy 1
Modal Structure, but other factors such as the cooling rate during commercial
production may
influence the formation of secondary phases such as martensite with varying
volume fraction.
Table 11 X-ray Diffraction Data for Alloy 1 After Solidification
(Modal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.583A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .876 A
Structure: Tetragonal
Ma rtensite Space group #: 139 (I4/mmm)
LP: a = .898 A
c = 3.018 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.093 A
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Deformation of the Alloy 1 with the Modal Structure (Structure #1, FIG. 1A) at
elevated
temperature induces homogenization and refinement of Modal Structure. Hot
rolling was
applied in this case but other processes including but not limited to hot
pressing, hot forging,
hot extrusion can achieve the similar effect. During hot rolling, the
dendrites in the Modal
5 Structure are broken up and refined, leading initially to the Homogenized
Modal Structure
(Structure #1a, FIG. 1A) formation. The refinement during the hot rolling
occurs through the
Nanophase Refinement (Mechanism #1, FIG. 1A) along with dynamic
recrystallization. The
Homogenized Modal Structure can be progressively refined by applying the hot
rolling
repetitively, leading to the Nanomodal Structure (Structure #2, FIG. 1A)
formation. FIG. 8a
10 shows the backscattered SEM micrograph of Alloy 1 after being hot rolled
from 50 mm to
¨1.7 mm at 1250 C. It can be seen that blocks of tens of microns in size are
resulted from the
dynamic recrystallization during the hot rolling, and the interior of the
grains is relatively
smooth indicating less amount of defects. TEM further reveals that sub-grains
of less than
several hundred nanometers in size are formed, as shown in FIG. 8b. X-ray
diffraction
15 analysis shows that the Nanomodal Structure of the Alloy 1 after hot
rolling contains mainly
austenite, with other phases such as ferrite and the iron manganese compound
as shown in
FIG. 9 and Table 12.
Table 12 X-ray Diffraction Data for Alloy 1 After Hot Rolling
20 (Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.595 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .896 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.113 A
Further deformation at ambient temperature (i.e., cold deformation) of the
Alloy 1 with the
Nanomodal Structure causes transformation into High Strength Nanomodal
Structure
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(Structure #3, FIG. 1A) through the Dynamic Nanophase Strengthening (Mechanism
#2, FIG.
1A). The cold deformation can be achieved by cold rolling and, tensile
deformation, or other
type of deformation such as punching, extrusion, stamping, etc. During the
cold deformation,
depending on alloy chemistries, a large portion of austenite in the Nanomodal
Structure is
transformed to ferrite with grain refinement. FIG. 10a shows the backscattered
SEM
micrograph of cold rolled Alloy 1. Compared to the smooth grains in the
Nanomodal
Structure after hot rolling, the cold deformed grains are rough indicating
severe plastic
deformation within the grains. Depending on alloy chemistry, deformation twins
can be
produced in some alloys especially by cold rolling, as displayed in FIG. 10a.
FIG. 10b
shows the TEM micrograph of the microstructure in cold rolled Alloy 1. It can
be seen that
in addition to dislocations generated by the deformation, refined grains due
to phase
transformation can also be found. The banded structure is related to the
deformation twins
caused by the cold rolling, corresponding to these in FIG 10a. X-ray
diffraction shows that
the High Strength Nanomodal Structure of the Alloy 1 after cold rolling
contains a significant
amount of ferrite phase in addition to the retained austenite and the iron
manganese
compound as shown in FIG. 11 and Table.
Table 13 X-ray Diffraction Data for Alloy 1 after Cold Rolling
(High Strength Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.588 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .871 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.10 A
Recrystallization occurs upon heat treatment of the cold deformed Alloy 1 with
High
Strength Nanomodal Structure (Structure #3,FIG. 1A and 1B) that transforms
into
Recrystallized Modal Structure (Structure #4,FIG. 1B). The TEM images of the
Alloy 1 after
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annealing are shown in ,FIG. 12. As it can be seen, equiaxed grains with sharp
and straight
boundaries are present in the structure and the grains are free of
dislocations, which is
characteristic feature of recrystallization. Depending on the annealing
temperature, the size
of recrystallized grains can range from 0.5 to 50 um. In addition, as shown in
electron
diffraction shows that austenite is the dominant phase after
recrystallization. Annealing twins
are occasionally found in the grains, but stacking faults are most often seen.
The formation
of stacking faults shown in the TEM image is typical for face-centered-cubic
crystal structure
of austenite. Backscattered SEM micrographs in FIG. 13 show the equiaxed
recrystallized
grains with the size of less than 10 um, consistent with TEM. The different
contrast of grains
(dark or bright) seen on SEM images suggests that the crystal orientation of
the grains is
random, since the contrast in this case is mainly originated from the grain
orientation. As a
result, any texture formed by the previous cold deformation is eliminated. X-
ray diffraction
shows that the Recrystallized Modal Structure of the Alloy 1 after annealing
contains
primarily austenite phase, with a small amount of ferrite and the iron
manganese compound
as shown in FIG. 14 and Table 14.
Table 14 X-ray Diffraction Data for Alloy 1 After Annealing
(Recrystallized Modal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.597 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .884 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.103 A
When the Alloy 1 with Recrystallized Modal Structure (Structure #4, FIG. 1B)
is subjected to
deformation at ambient temperature, Nanophase Refinement & Strengthening
(Mechanism
#4, FIG. 1B) is activated leading to formation of the Refined High Strength
Nanomodal
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Structure (Structure #5, FIG. 1B). In this case, deformation was a result of
tensile testing and
gage section of the tensile sample after testing was analyzed. FIG. 15 shows
the bright-field
TEM micrographs of the microstructure in the deformed Alloy 1. Compared to the
matrix
grains that were initially almost dislocation-free in the Recrystallized Modal
Structure after
annealing, the application of stress generates a high density of dislocations
within the matrix
grains. At the end of tensile deformation (with a tensile elongation greater
than 50%),
accumulation of large number of dislocations is observed in the matrix grains.
As shown in
FIG. 15a, in some areas (for example the area at the lower part of the FIG.
15a), dislocations
form a cell structure and the matrix remains austenitic. In other areas, where
the dislocation
density is sufficiently high, transformation is induced from austenite to
ferrite (for example
the upper and right part of the FIG. 15a) that results in substantial
structure refinement. FIG.
15b shows local "pocket" of the transformed refined microstructure and
selected area electron
diffraction pattern corresponds to ferrite. Structural transformation into
Refined High
Strength Nanomodal Structure (Structure #5, FIG. 1B) in the randomly
distributed "pockets"
is a characteristic feature of the steel alloys herein. FIG. 16 shows the
backscattered SEM
images of the Refined High Strength Nanomodal Structure. Compared to the
Recrystallized
Modal Structure, the boundaries of matrix grains become less apparent, and the
matrix is
obviously deformed. Although the details of deformed grains cannot be revealed
by SEM,
the change caused by the deformation is enormous compared to the
Recrystallized Modal
Structure that was demonstrated in TEM images. X-ray diffraction shows that
the Refined
High Strength Nanomodal Structure of the Alloy 1 after tensile deformation
contains a
significant amount of ferrite and austenite phases. Very broad peaks of
ferrite phase (alpha-
Fe) are seen in the XRD pattern, suggesting significant refinement of the
phase. The iron
manganese compound is also present. Additionally, a hexagonal phase with space
group
#186 (P63me) was identified in the gage section of the tensile sample as shown
in FIG. 17 and
Table 15.
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Table 15 X-ray Diffraction Data for Alloy 1 After Tensile Deformation
(Refined High Strength Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
7- Fe
Space group #: 5 (Fm3m)
LP: a = 3.586 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .873 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.159 A
Structure: Hexagonal
Hexagonal phase 1
Space group #: 186 (P63mc)
LP: a = 3.013 A, c = 6.183 A
This Case Example demonstrates that alloys listed in Table 2 including Alloy 1
exhibit a
structural development pathway with novel enabling mechanisms illustrated in
FIGS. 1A and
1B leading to unique microstructures with nanoscale features.
Case Example #2 Structural Development Pathway in Alloy 2
Laboratory slab with thickness of 50 mm was cast from Alloy 2 that was then
laboratory
processed by hot rolling, cold rolling and annealing at 850 C for 10 mm as
described in Main
Body section of current application. Microstructure of the alloy was examined
at each step of
processing by SEM, TEM and x-ray analysis.
For SEM study, the cross section of the slab samples was ground on SiC
abrasive papers with
reduced grit size, and then polished progressively with diamond media paste
down to 1 um.
The final polishing was done with 0.02 um grit 5i02 solution. Microstructures
were
examined by SEM using an EVO-MA10 scanning electron microscope manufactured by
Carl
Zeiss SMT Inc. To prepare TEM specimens, the samples were first cut with EDM,
and then
thinned by grinding with pads of reduced grit size every time. Further
thinning to make foils
to ¨60 um thickness was done by polishing with 9 um, 3 um and 1 um diamond
suspension
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solution respectively. Discs of 3 mm in diameter were punched from the foils
and the final
polishing was fulfilled with electropolishing using a twin-jet polisher. The
chemical solution
used was a 30% nitric acid mixed in methanol base. In case of insufficient
thin area for TEM
observation, the TEM specimens may be ion-milled using a Gatan Precision Ion
Polishing
5 System (PIPS). The ion-milling usually is done at 4.5 keV, and the
inclination angle is
reduced from 4 to 2 to open up the thin area. The TEM studies were done
using a JEOL
2100 high-resolution microscope operated at 200 kV. X-ray diffraction was done
using a
Panalytical X'Pert MPD diffractometer with a Cu Koc x-ray tube and operated at
45 kV with a
filament current of 40 mA. Scans were run with a step size of 0.01 and from
25 to 95 two-
10 theta with silicon incorporated to adjust for instrument zero angle
shift. The resulting scans
were then subsequently analyzed using Rietveld analysis using Siroquant
software.
Modal Structure (Structure #1, FIG. 1A) is formed in Alloy 2 slab cast at 50
mm thick, which
is characterized by dendritic structure. Due to the presence of a boride phase
(M2B), the
dendritic structure is more evident than in Alloy 1 where borides are absent.
FIG. 18a shows
15 .. the backscattered SEM of Modal Structure that exhibits a dendritic
matrix (in bright contrast)
with borides at the boundary (in dark contrast). TEM studies show that the
matrix phase is
composed of austenite (gamma-Fe) with stacking faults (FIG. 18b). Similar to
Alloy 1, the
presence of stacking faults indicates the matrix phase is austenite. Also
shown in TEM is the
boride phase that appears dark in. FIG. 18b at the boundary of austenite
matrix phase. X-ray
20 diffraction analysis data in. FIG. 19 and Table 16 shows that the Modal
Structure contains
austenite, M2B, ferrite, and iron manganese compound. Similar to Alloy 1,
austenite is the
dominant phase in the Alloy 2 Modal Structure, but other phases may be present
depending
on alloy chemistry.
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Table 16 X-ray Diffraction Data for Alloy 2 After Solidification
(Modal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.577 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .850 A
Structure: Tetragonal
M2B
Space group #: 140 (I4/mcm)
LP: a = 5.115 A , c = 4. 6 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.116 A
Following the flowchart in FIG. 1A, deformation of the Alloy 2 with the Modal
Structure
(Structure #1, FIG. 1A) at elevated temperature induces homogenization and
refinement of
Modal Structure. Hot rolling was applied in this case but other processes
including but not
limited to hot pressing, hot forging, hot extrusion can achieve a similar
effect. During the hot
rolling, the dendrites in the Modal Structure are broken up and refined,
leading initially to the
Homogenized Modal Structure (Structure #1a, FIG. 1A) formation. The refinement
during
the hot rolling occurs through the Nanophase Refinement (Mechanism #1, FIG.
1A) along
with dynamic recrystallization. The Homogenized Modal Structure can be
progressively
refined by applying the hot rolling repetitively, leading to the Nanomodal
Structure (Structure
#2, FIG. 1A) formation. FIG. 20a shows the backscattered SEM micrograph of hot
rolled
Alloy 2. Similar to Alloy 1, the dendritic Modal Structure is homogenized
while the boride
phase is randomly distributed in the matrix. TEM shows that the matrix phase
is partially
recrystallized as a result of dynamic recrystallization during hot rolling, as
shown in FIG.
20b. The matrix grains are on the order of 500 nm, which is finer than in
Alloy 1 due to the
pinning effect of borides. X-ray diffraction analysis shows that the Nanomodal
Structure of
Alloy 2 after hot rolling contains mainly austenite phase and M2B, with other
phases such as
ferrite and iron manganese compound as shown in FIG. 21 and Table 17.
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Table 17 X-ray Diffraction Data for Alloy 2 After Hot Rolling
(Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
7- Fe
Space group #: 5 (Fm3m)
LP: a = 3.598 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .853 A
Structure: Tetragonal
M2B
Space group #: 140 (14/m cm)
LP: a = 5.1' 3 A , c = 4.18 A
Structure: Cubic
Iron manganese compound
Space group #: 5 (Fm3m)
LP: a = 4.180 A
Deformation of the Alloy 2 with the Nanomodal Structure but at ambient
temperature (i.e.,
cold deformation) leads to formation of High Strength Nanomodal Structure
(Structure #3,
FIG. 1A) through the Dynamic Nanophase Strengthening (Mechanism #2, FIG. 1A).
The
cold deformation can be achieved by cold rolling, tensile deformation, or
other type of
deformation such as punching, extrusion, stamping, etc. Similarly in Alloy 2
during cold
deformation, a great portion of austenite in the Nanomodal Structure is
transformed to ferrite
with grain refinement. FIG. 22a shows the backscattered SEM micrograph of the
microstructure in the cold rolled Alloy 2. Deformation is concentrated in the
matrix phase
around the boride phase. FIG. 22b shows the TEM micrograph of the cold rolled
Alloy 2.
Refined grains can be found due to the phase transformation. Although
deformation twins
are less evident in SEM image, TEM shows that they are generated after the
cold rolling,
similar to Alloy 1. X-ray diffraction shows that the High Strength Nanomodal
Structure of
the Alloy 2 after cold rolling contains a significant amount of ferrite phase
in addition to the
M2B, retained austenite and a new hexagonal phase with space group #186
(P63õ,) as shown
in FIG. 23 and Table 18.
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Table 18 X-ray Diffraction Data for Alloy 2 After Cold Rolling
(High Strength Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
7- Fe
Space group #: 5 (Fm3m)
LP: a = 3.551 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .874 A
Structure: Tetragonal
M2B
Space group #: 140 (I4/mcm)
LP: a = 5.1' 5 A , c = 4. 03 A
Structure: Hexagonal
Hexagonal phase
Space group #: 186 (P63mc)
LP: a = .96 A, c = 6. 7' A
Recrystallization occurs upon annealing of the cold deformed Alloy 2 with High
Strength
Nanomodal Structure (Structure #3, FIG. 1A and 1B) that transforms into
Recrystallized
Modal Structure (Structure #4, FIG. 1B). The recrystallized microstructure of
the Alloy 2
after annealing is shown by TEM images in FIG. 24. As it can be seen, equiaxed
grains with
sharp and straight boundaries are present in the structure and the grains are
free of
dislocations, which is a characteristic feature of recrystallization. The size
of recrystallized
grains is generally less than 5 um due to the pinning effect of boride phase,
but larger grains
are possible at higher annealing temperatures. Moreover, electron diffraction
shows that
austenite is the dominant phase after recrystallization and stacking faults
are present in the
austenite, as shown in FIG. 24b. The formation of stacking faults also
indicates formation of
face-centered-cubic austenite phase. Backscattered SEM micrographs in FIG. 25
show the
equiaxed recrystallized grains with the size of less than 5 um, with boride
phase randomly
distributed. The different contrast of grains (dark or bright) seen on SEM
images suggests
that the crystal orientation of the grains is random, since the contrast in
this case is mainly
originated from the grain orientation. As a result, any texture formed by the
previous cold
deformation is eliminated. X-ray diffraction shows that the Recrystallized
Modal Structure
of the Alloy 2 after annealing contains primarily austenite phase, with M2B, a
small amount
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of ferrite, and a hexagonal phase with space group #186 (P63õ,) as shown in
FIG. 26 and
Table 19.
Table 19 X-ray Diffraction Data for Alloy 2 After Annealing
(Recrystallized Modal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.597 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .878 A
Structure: Tetragonal
M2B
Space group #: 140 (14/m cm)
LP: a = 5.153 A , c = 4.170 A
Structure: Hexagonal
Hexagonal phase
Space group #: 186 (P63mc)
LP: a = .965 A, c = 6. 70 A
Deformation of Recrystallized Modal Structure (Structure #4, FIG. 1B) leads to
formation of
the Refined High Strength Nanomodal Structure (Structure #5, FIG. 1B) through
Nanophase
Refinement & Strengthening (Mechanism #4, FIG. 1B). In this case, deformation
was a
result of tensile testing and the gage section of the tensile sample after
testing was analyzed.
FIG. 27 shows the micrographs of microstructure in the deformed Alloy 2.
Similar to Alloy
1, the initially dislocation-free matrix grains in the Recrystallized Modal
Structure after
annealing are filled with a high density of dislocations upon the application
of stress, and the
accumulation of dislocations in some grains activates the phase transformation
from austenite
to ferrite, leading to substantial refinement. As shown in FIG. 27a, refined
grains of 100 to
300 nm in size are shown in a local "pocket" where transformation occurred
from austenite to
ferrite. Structural transformation into Refined High Strength Nanomodal
Structure (Structure
#5, FIG 1B) in the "pockets" of matrix grains is a characteristic feature of
the steel alloys
herein. FIG. 27b shows the backscattered SEM images of the Refined High
Strength
Nanomodal Structure. Similarly, the boundaries of matrix grains become less
apparent after
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the matrix is deformed. X-ray diffraction shows that a significant amount of
austenite
transformed to ferrite although the four phases remain as in the
Recrystallized Modal
Structure. The transformation resulted in formation of Refined High Strength
Nanomodal
Structure of the Alloy 2 after tensile deformation. Very broad peaks of
ferrite phase (a-Fe)
5 .. are seen in the XRD pattern, suggesting significant refinement of the
phase. As in Alloy 1, a
new hexagonal phase with space group #186 (P63,Tie) was identified in the gage
section of the
tensile sample as shown in FIG. 28 and Table 20.
Table 20 X-ray Diffraction Data for Alloy 2 After Tensile Deformation
10 (Refined High Strength Nanomodal Structure)
Phases Identified Phase Details
Structure: Cubic
y - Fe
Space group #: 5 (Fm3m)
LP: a = 3.597 A
Structure: Cubic
a - Fe
Space group #: 9 (Im3m)
LP: a = .898 A
Structure: Tetragonal
M2B
Space group #: 140 (I4/mcm)
LP: a = 5.149 A, c = 4.181 A
Structure: Hexagonal
Hexagonal phase
Space group #: 186 (P63mc)
LP: a = .961 A, c = 6. 71 A
This Case Example demonstrates that alloys listed in Table 2 including Alloy 2
exhibit a
structural development pathway with the mechanisms illustrated in FIGS. 1A and
1B leading
to unique microstructures with nanoscale features.
Case Example #3 Tensile Properties at Each Step of Processing
Slabs with thickness of 50 mm were laboratory cast from the alloys listed in
Table 21
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 mm as described in Main Body
section of current
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application. Tensile properties were measured at each step of processing on an
Instron 3369
mechanical testing frame using Instron's Bluehill control software. All tests
were conducted
at room temperature, with the bottom grip fixed and the top grip set to travel
upwards at a
rate of 0.012 mm/s. Strain data was collected using Instron's Advanced Video
Extensometer.
Alloys were weighed out into charges ranging from 3,000 to 3,400 grams using
commercially
available ferroadditive powders with known chemistry and impurity content
according to the
atomic ratios in Table 2. Charges were loaded into zirconia coated silica
crucibles which
were placed into an Indutherm VTC800V vacuum tilt casting machine. The machine
then
evacuated the casting and melting chambers and backfilled with argon to
atmospheric
pressure several times prior to casting to prevent oxidation of the melt. The
melt was heated
with a 14 kHz RF induction coil until fully molten, approximately 5.25 to 6.5
minutes
depending on the alloy composition and charge mass. After the last solids were
observed to
melt it was allowed to heat for an additional 30 to 45 seconds to provide
superheat and ensure
melt homogeneity. The casting machine then evacuated the melting and casting
chambers
and tilted the crucible and poured the melt into a 50 mm thick, 75 to 80 mm
wide, and 125
mm deep channel in a water cooled copper die. The melt was allowed to cool
under vacuum
for 200 seconds before the chamber was filled with argon to atmospheric
pressure. Tensile
specimens were cut from as-cast slabs by wire EDM and tested in tension.
Results of tensile
testing are shown in Table 21. As it can be seen, ultimate tensile strength of
the alloys herein
in as-cast condition varies from 411 to 907 MPa. The tensile elongation varies
from 3.7 to
24.4%. Yield strength is measured in a range from 144 to 514 MPa.
Prior to hot rolling, laboratory cast slabs were loaded into a Lucifer EHS3GT-
B18 furnace to
heat. The furnace set point varies between 1000 C to 1250 C depending on alloy
melting
point. The slabs were allowed to soak for 40 minutes prior to hot rolling to
ensure they reach
the target temperature. Between hot rolling passes the slabs are returned to
the furnace for 4
minutes to allow the slabs to reheat. Pre-heated slabs were pushed out of the
tunnel furnace
into a Fenn Model 061 2 high rolling mill. The 50 mm casts are hot rolled for
5 to 8 passes
through the mill before being allowed to air cool defined as first campaign of
hot rolling.
After this campaign the slab thickness was reduced between 80.4 to 87.4%.
After cooling,
the resultant sheet samples were sectioned to 190 mm in length. These sections
were hot
rolled for an additional 3 passes through the mill with reduction between 73.1
to 79.9% to a
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final thickness of between 2.1 and 1.6 mm. Detailed information on hot rolling
conditions for
each alloy herein is provided in Table 22. Tensile specimens were cut from hot
rolled sheets
by wire EDM and tested in tension. Results of tensile testing are shown in
Table 22. After
hot rolling, ultimate tensile strength of the alloys herein varies from 921 to
1413 MPa. The
tensile elongation varies from 12.0 to 77.7%. Yield strength is measured in a
range from 264
to 574 MPa. See, Structure 2 in FIG. 1A.
After hot rolling, resultant sheets were media blasted with aluminum oxide to
remove the mill
scale and were then cold rolled on a Fenn Model 061 2 high rolling mill. Cold
rolling takes
multiple passes to reduce the thickness of the sheet to targeted thickness,
generally 1.2 mm.
Hot rolled sheets were fed into the mill at steadily decreasing roll gaps
until the minimum gap
is reached. If the material has not yet hit the gauge target, additional
passes at the minimum
gap were used until the targeted thickness was reached. Cold rolling
conditions with the
number of passes for each alloy herein are listed in Table 23. Tensile
specimens were cut
from cold rolled sheets by wire EDM and tested in tension. Results of tensile
testing are
shown in Table 23. Cold rolling leads to significant strengthening with
ultimate tensile
strength in the range from 1356 to 1831 MPa. The tensile elongation of the
alloys herein in
cold rolled state varies from 1.6 to 32.1%. Yield strength is measured in a
range from 793 to
1645 MPa. It is anticipated that higher ultimate tensile strength and yield
strength can be
achieved in alloys herein by larger cold rolling reduction (>40%) that in our
case is limited by
laboratory mill capability. With more rolling force, it is anticipated that
ultimate tensile
strength could be increased to at least 2000 MPa and yield strength to at
least 1800 MPa.
Tensile specimens were cut from cold rolled sheet samples by wire EDM and
annealed at
850 C for 10 mm in a Lucifer 7HT-K12 box furnace. Samples were removed from
the
furnace at the end of the cycle and allowed to cool to room temperature in
air. Results of
tensile testing are shown in Table 24. As it can be seen, recrystallization
during annealing of
the alloys herein results in property combinations with ultimate tensile
strength in the range
from 939 to 1424 MPa and tensile elongation from 15.8 to 77.0%. Yield strength
is
measured in a range from 420 to 574 MPa. FIG. 29 to FIG. 31 represent plotted
data at each
processing step for Alloy 1, Alloy 13, and Alloy 17, respectively.
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Table 21 Tensile Properties of Alloys in As-Cast State
Yield Ultimate Tensile
Tensile Elongation
Alloy Strength Strength
(%)
(MPa) (MPa)
289 527 10.4
Alloy 1 288 548 9.3
260 494 8.4
244 539 10.4
Alloy 2 251 592 11.6
249 602 13.1
144 459 4.6
Alloy 13 156 411 4.5
163 471 5.7
223 562 24.4
Alloy 17 234 554 20.7
235 585 23.3
396 765 8.3
Alloy 24 362 662 5.7
404 704 7.0
282 668 5.1
Alloy 25 329 753 5.0
288 731 5.5
471 788 4.1
Alloy 25 514 907 6.0
483 815 3.7
277 771 3.7
Alloy 27 278 900 4.9
267 798 4.5
152 572 11.1
Alloy 34 168 519 11.6
187 545 12.9
164 566 15.9
Alloy 35 172 618 16.6
162 569 16.4
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Table 22 Tensile Properties of Alloys in Hot Rolled State
Ultimate
First Second Yield Tensile
Tensile
Alloy Condition Campaign Campaign Strength
Strength Elongation
Reduction Reduction (MPa) (%)
(MPa)
273 1217 50.0
Hot Rolled 80.5%, 75.1%,
Alloy 1 264 1216 52.1
95.2% 6 Passes 3 Passes
285 1238 52.7
480 1236 45.3
Hot Rolled 87.4%, 73.1%,
Alloy 2 454 1277 41.9
96.6% 7 Passes 3 Passes
459 1219 48.2
287 1116 18.8
Hot Rolled 81.1%, 79.8%,
Alloy 13 274 921 15.3
96.0% 6 Passes 3 Passes
293 1081 19.3
392 947 73.3
Hot Rolled 81.2%, 79.1%,
Alloy 17 363 949 74.8
96.1% 6 Passes 3 Passes
383 944 77.7
519 1176 21.4
Hot Rolled, 81.1%, 79.9%,
Alloy 24 521 1088 18.2
96.2% 6 Passes 3 Passes
508 1086 17.9
502 1105 12.4
Hot Rolled 81.0%, 79.4%,
Alloy 25 524 1100 12.3
96.1% 6 Passes 3 Passes
574 1077 12.0
508 1401 20.9
Hot Rolled, 80.4%, 78.9%,
Alloy 27 534 1405 22.4
95.9% 6 Passes 3 Passes
529 1413 19.7
Alloy 34 Hot Rolled, 80.7%, 6 80.1 %, 3 346 1188
56.5
96.2% Passes Passes 323 1248 58.7
303 1230 53.4
Alloy 35 Hot Rolled, 80.8%, 6 79.9%, 3 327 1178
63.3
96.1% Passes Passes 317 1170 61.2
305 1215 59.6
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Table 23 Tensile Properties of Alloys in Cold Rolled State
Yield Strength Ultimate Tensile
Tensile Elongation
Alloy Condition
(MPa) Strength (MPa) (%)
Cold Rolled 798 1492 28.5
20.3%,
4 Passes 793 1482 32.1
Alloy 1 1109 1712 21.4
Cold Rolled
39.6%, 1142 1726 23.0
29 Passes 1203 1729 21.2
Cold Rolled 966 1613 13.4
28.5%, 998 1615 15.4
5 Passes
Alloy 2 1053 1611 20.6
Cold Rolled 1122 1735 20.3
39.1%,
1270 1744 18.3
19 passes
Cold Rolled 1511 1824 9.5
Alloy 13 36.0%, 1424 1803 7.7
24 Passes 1361 1763 5.1
Cold Rolled 1020 1357 24.2
Alloy 17 38.5%, 1007 1356 24.9
8 Passes 1071 1357 24.9
1363 1584 1.9
Cold Rolled
Alloy 24 38.2%, 1295 1601 2.5
23 Passes 1299 1599 3.0
1619 1761 1.9
Cold Rolled
Alloy 25 38.0%, 1634 1741 1.7
42 Passes 1540 1749 1.6
Cold Rolled 1632 1802 2.7
Alloy 27 39.4%, 1431 1804 4.1
40 Passes 1645 1831 4.1
1099 1640 14.7
Cold Rolled
Alloy 34 35.%, 14 840 1636 17.5
Passes 1021 1661 18.5
Cold Rolled 996 1617 23.8
Alloy 35 35.5%, 12 1012 1614 24.5
Passes 1020 1616 23.3
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Table 24 Tensile Properties of Alloys in Annealed State
All Yield Strength Ultimate Tensile Tensile Elongation
oy
(MPa) Strength (MPa) (%)
436 1221 54.9
Alloy 1 443 1217 56.0
431 1216 59.7
438 1232 49.7
431 1228 49.8
431 1231 49.4
Alloy 2
484 1278 48.3
485 1264 45.5
479 1261 48.7
441 1424 41.7
Alloy 13 440 1412 41.4
429 1417 42.7
420 946 74.6
Alloy 17 421 939 77.0
425 961 74.9
554 1151 23.5
Alloy 24 538 1142 24.3
562 1151 24.3
500 1274 16.0
Alloy 25 502 1271 15.8
483 1280 16.3
538 1385 20.6
Alloy 27 574 1397 20.9
544 1388 21.8
467 1227 56.7
Alloy 27 476 1232 52.7
462 1217 51.6
439 1166 56.3
Alloy 27 438 1166 59.0
440 1177 58.3
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This Case Example demonstrates that due to the unique mechanisms and
structural pathway
shown in FIGS. 1A and 1B, the structures and resulting properties in steel
alloys herein can
vary widely leading to the development of 3rd Generation AHSS.
Case Example #4 Cyclic Reversibility During Cold Rolling and Recrystallization
Slabs with thickness of 50 mm were laboratory cast from Alloy 1 and Alloy 2
according to
the atomic ratios provided in Table 2 and hot rolled into sheets with final
thickness of 2.31
mm for Alloy 1 sheet and 2.35 mm for Alloy 2 sheet. Casting and hot rolling
procedures are
described in Main Body section of current application. Resultant hot rolled
sheet from each
alloy was used for demonstration of cyclic structure/property reversibility
through cold
rolling/annealing cycles.
Hot rolled sheet from each alloy was subjected to three cycles of cold rolling
and annealing.
Sheet thicknesses before and after hot rolling and cold rolling reduction at
each cycle are
listed in Table 25. Annealing at 850 C for 10 min was applied after each cold
rolling.
Tensile specimens were cut from the sheet in the initial hot rolled state and
at each step of the
cycling. Tensile properties were measured on an Instron 3369 mechanical
testing frame
using Instron's Bluehill control software. All tests were conducted at room
temperature, with
the bottom grip fixed and the top grip set to travel upwards at a rate of
0.012 mm/s. Strain
data was collected using Instron's Advanced Video Extensometer.
The results of tensile testing are plotted in FIG. 32 for Alloy 1 and Alloy 2
showing that cold
rolling results in significant strengthening of both alloys at each cycle with
average ultimate
tensile strength of 1500 MPa in Alloy 1 and 1580 MPa in Alloy 2. Both cold
rolled alloys
show a loss in ductility as compared to the hot rolled state. However,
annealing after cold
rolling at each cycle results in tensile property recovery to the same level
with high ductility.
Tensile properties for each tested sample are listed in Table 26 and Table 27
for Alloy 1 and
Alloy 2, respectively. As it can be seen, Alloy 1 has ultimate tensile
strength from 1216 to
1238 MPa in hot rolled state with ductility from 50.0 to 52.7% and yield
strength from 264 to
285 MPa. In cold rolled state, the ultimate tensile strength was measured in
the range from
1482 to 1517 MPa at each cycle. Ductility was found consistently in the range
from 28.5 to
32.8% with significantly higher yield strength of 718 to 830 MPa as compared
to that in hot
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rolled condition. Annealing at each cycle resulted in restoration of the
ductility to the range
from 47.7 to 59.7% with ultimate tensile strength from 1216 to 1270 MPa. Yield
strength
after cold rolling and annealing is lower than that after cold rolling and was
measured in the
range from 431 to 515 MPa that is however higher than that in initial hot
rolled condition.
Similar results with property reversibility between cold rolled and annealed
material through
cycling were observed for Alloy 2 (FIG. 32b). In initial hot rolled state,
Alloy 2 has ultimate
tensile strength from 1219 to 1277 MPa with ductility from 41.9 to 48.2% and
yield strength
from 454 to 480 MPa. Cold rolling at each cycle results in the material
strengthening to the
ultimate tensile strength from 1553 to 1598 MPa with ductility reduction to
the range from
20.3 to 24.1%. Yield strength was measured from 912 to 1126 MPa. After
annealing at each
cycle, Alloy 2 has ultimate tensile strength from 1231 to 1281 MPa with
ductility from 46.9
to 53.5%. Yield strength in Alloy 2 after cold rolling and annealing at each
cycle is similar to
that in hot rolled condition and varies from 454 to 521 MPa.
Table 25 Sample Thickness and Cycle Reduction at Cold Rolling Steps
Allo Rolling Initial Thickness Final Thickness
Cycle Reduction
y
Cycle (mm) (mm) (%)
1 2.35 1.74 26.0
Alloy 1 2 1.74 1.32 24.1
3 1.32 1.02 22.7
1 2.31 1.85 19.9
Alloy 2 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
1st Cycle 2nd Cycle 3rd Cycle
Property Hot Rolled Cold Cold Cold
Annealed Annealed Annealed
Rolled Rolled Rolled
Ultimate 1217 1492 1221 1497 1239 1517 1270
Tensile
1216 1482 1217 1507 1269 1507 1262
Strength
(MPa) 1238 1216 1503 1260 1507 1253
Yield 273 798 436 775 487 820 508
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Strength 264 793 443 718 466 796 501
(MPa)
285 * 431 830 488 809 515
Tensile 50.0 28.5 54.9 32.8 57.5 32.1 50.5
Elongation 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
* Specimens slipped in the grips / data is not available
Table 27 Tensile Properties of Alloy 2 Through Cold Rolling/Annealing Cycles
1st Cycle 2nd Cycle 3rd Cycle
Property Hot Rolled Cold Cold Cold
Annealed Annealed
Annealed
Rolled Rolled Rolled
Ultimate 1236 1579 1250 1553 1243 1596 1231
Tensile
1277 * 1270 1568 1255 1589 1281
Strength
(MPa) 1219 * 1240 1566 1242 1598 1269
Yield 480 1126 466 983 481 1006 475
Strength 454 * 468 969 521 978 507
(MPa) 459 * 454 912 497 1011 518
Tensile 45.3 20.3 53.0 24.1 51.1 22.3 46.9
Elongation 41.9 * 51.2 23.1 52.3 23.2 53.5
(%) 48.2 * 51.1 21.6 49.9 21.0 47.9
* Specimens slipped in the grips / data is not available
This Case Example demonstrates that the High Strength Nanomodal Structure
(Structure #3,
FIG. 1A) that forms in the alloys listed in Table 2 after cold rolling can be
recrystallized by
applying an anneal to produce a Recrystallized Modal Structure (Structure #4,
FIG. 1B).
This structure can be further deformed through cold rolling or other cold
deformation
approaches to undergo Nanophase Refinement and Strengthening (Mechanism #4,
FIG. 1B)
leading to formation of the Refined High Strength Nanomodal Structure
(Structure #5, FIG.
1B). The Refined High Strength Nanomodal Structure (Structure #5, FIG. 1B) can
in turn be
recrystallized and the process can be started over with full
structure/property reversibility
through multiple cycles. The ability for the mechanisms to be reversible
enables the
production of finer gauges which are important for weight reduction when using
AHSS as
well as property recovery after any damage caused by deformation.
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Case Example #5 Bending Ability
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 28
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
5 cold rolling and annealing at 850 C for 10 mm as described in Main Body
section of current
application. Resultant sheet from each alloy with final thickness of -1.2 mm
and
Recrystallized Modal Structure (Structure #4, FIG. 1B) was used to evaluate
bending
response of alloys herein.
Bend tests were performed using an Instron 5984 tensile test platform with an
Instron W-
10 6810 guided bend test fixture according to specifications outlined in
the ISO 7438
International Standard Metallic materials¨Bend test (International
Organization for
Standardization, 2005). Test specimens were cut by wire EDM to a dimension of
20 mm x
55 mm x sheet thickness. No special edge preparation was done to the samples.
Bend tests
were performed using an Instron 5984 tensile test platform with an Instron W-
6810 guided
15 bend test fixture. Bend tests were performed according to specifications
outlined in the ISO
7438 International Standard Metallic materials¨Bend test (International
Organization for
Standardization, 2005).
The test was performed by placing the test specimen on the fixture supports
and pushing with
a former as shown in FIG. 33.
20 The distance between supports, 1, was fixed according to ISO 7438 during
the test at:
1= (D +3a) + -a
Equation 1
- 2
Prior to bending, the specimens were lubricated on both sides with 3 in 1 oil
to reduce friction
with the test fixture. This test was performed with a 1 mm diameter former.
The former was
pushed downward in the middle of the supports to different angles up to 180
or until a crack
25 appeared. The bending force was applied slowly to permit free plastic
flow of the material.
The displacement rate was calculated based on the span gap of each test in
order to have a
constant angular rate and applied accordingly.
Absence of cracks visible without the use of magnifying aids was considered
evidence that
the test piece withstood the bend test. If a crack was detected, the bend
angle was measured
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manually with a digital protractor at the bottom of the bend. The test
specimen was then
removed from the fixture and examined for cracking on the outside of the bend
radius. The
onset of cracking could not be conclusively determined from the force-
displacement curves
and was instead easily determined by direct observation with illumination from
a flashlight.
Results of the bending response of the alloys herein are listed in Table 28
including initial
sheet thickness, former radius to sheet thickness ratio (r/t) and maximum bend
angle before
cracking. All alloys listed in the Table 28 did not show cracks at 90 bend
angle. The
majority of the alloys herein have capability to be bent at 1800 angle without
cracking.
Example of the samples from Alloy 1 after bend testing to 180 is shown in
FIG. 34.
Table 7 Bend Test Results for Selected Alloys
Former
Thickness Maximum Bend
Alloy Diameter r/t
(mm) (mm) Angle ( )
1.185 0.401 180
1.200 0.396 180
1.213 0.392 180
1.223 0.388 180
Alloy 1 0.95
1.181 0.402 180
1.187 0.400 180
1.189 0.399 180
1.206 0.394 180
1.225 0.388 180
1.230 0.386 180
1.215 0.391 180
1.215 0.391 180
Alloy 2 0.95
1.215 0.391 180
1.224 0.388 180
1.208 0.393 180
1.208 0.393 180
1.212 0.392 180
Alloy 3 0.95 1.186 0.401 180
1.201 0.396 180
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Former
Thickness Maximum Bend
Alloy Diameter r/t
(mm) Angle ( )
(mm)
1.227 0.387 180
Alloy 4 0.95 1.185 0.401 180
1.187 0.400 180
1.199 0.396 110
Alloy 5 0.95
1.196 0.397 90
1.259 0.377 160
Alloy 6 0.95 1.202 0.395 165
1.206 0.394 142
1.237 0.384 104
Alloy 7 0.95
1.236 0.384 90
1.278 0.372 180
Alloy 9 0.95 1.197 0.397 180
1.191 0.399 180
1.226 0.387 180
1.208 0.393 100
Alloy 10 0.95
1.208 0.393 180
1.205 0.394 180
1.240 0.383 180
Alloy 11 0.95 1.214 0.391 180
1.205 0.394 180
1.244 0.382 180
Alloy 12 0.95 1.215 0.391 180
1.205 0.394 180
1.222 0.389 180
Alloy 13 0.95 1.191 0.399 180
1.188 0.400 180
1.239 0.383 180
Alloy 14 0.95 1.220 0.389 180
1.214 0.391 180
1.247 0.381 180
Alloy 15 0.95
1.224 0.388 180
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Former
Thickness Maximum Bend
Alloy Diameter r/t
(mm) Angle ( )
(mm)
1.224 0.388 180
1.244 0.382 180
Alloy 16 0.95 1.224 0.388 180
1.199 0.396 180
1.233 0.385 180
Alloy 17 0.95 1.213 0.392 180
1.203 0.395 180
1.222 0.389 160
Alloy 18 0.95
1.218 0.390 135
1.266 0.375 180
Alloy 19 0.95 1.243 0.382 180
1.242 0.382 180
1.242 0.382 180
Alloy 20 0.95 1.222 0.389 180
1.220 0.389 180
1.255 0.378 180
Alloy 21 0.95 1.228 0.387 180
1.229 0.386 180
1.240 0.383 180
Alloy 22 0.95 1.190 0.399 180
1.190 0.399 180
1.190 0.399 180
Alloy 23 0.95 1.199 0.396 180
1.193 0.398 180
1.222 0.389 180
Alloy 28 0.95 1.206 0.394 180
1.204 0.395 180
1.219 0.390 180
Alloy 29 0.95 1.217 0.390 180
1.206 0.394 180
Alloy 30 0.95 1.215 0.391 180
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Former
Thickness Maximum Bend
Alloy Diameter r/t
(mm) Angle ( )
(mm)
1.212 0.392 175
1.200 0.396 180
1.211 0.392 150
Alloy 31 0.95
1.209 0.393 131
1.222 0.389 180
Alloy 32 0.95 1.221 0.389 180
1.210 0.393 180
In order to be made into complex parts for automobile and other uses, an AHSS
needs to
exhibit both bulk sheet formability and edge sheet formability. This Case
Example
demonstrates good bulk sheet formability of the alloys in Table 2 through bend
testing.
Case Example #6 Punched Edge vs EDM Cut Tensile Properties
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table
2according to the atomic ratios provided in Table 2 and laboratory processed
by hot rolling,
cold rolling and annealing at 850 C for 10 min as described herein. Resultant
sheet from
each alloy with final thickness of 1.2 mm and Recrystallized Modal Structure
(Structure #4,
FIG. 1B) were used to evaluate the effect of edge damage on alloy properties
by cutting
tensile specimens by wire electrical discharge machining (wire-EDM) (which
represents the
control situation or relative lack of shearing and formation of an edge
without a compromise
in mechanical properties) and by punching (to identify a mechanical property
loss due to
shearing). It should be appreciated that shearing (imposition of a stress
coplanar with a
material cross-section) may occur herein by a number of processing options,
such as piercing,
perforating, cutting or cropping (cutting off of an end of a given metal
part).
Tensile specimens in the ASTM E8 geometry were prepared using both wire EDM
cutting
and punching. Tensile properties were measured on an Instron 5984 mechanical
testing
frame using Instron' s Bluehill control software. All tests were conducted at
room
temperature, with the bottom grip fixed and the top grip set to travel upwards
at a rate of
0.012 mm/s. Strain data was collected using Instron' s Advanced Video
Extensometer.
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Tensile data is shown in Table 29 and illustrated in FIG. 35a for selected
alloys. Decrease in
properties is observed for all alloys tested but the level of this decrease
varies significantly
depending on alloy chemistry. Table 30 summarizes a comparison of ductility in
punched
samples as compared to that in the wire EDM cut samples. In FIG. 35b
corresponding tensile
5 curves are shown for the selected alloy demonstrating mechanical behavior
as a function of
austenite stability. For selected alloys herein, austenite stability is
highest in Alloy 12 that
shows high ductility and lowest in Alloy 13 that shows high strength.
Correspondingly,
Alloy 12 demonstrated lowest loss in ductility in punched specimens vs EDM cut
(29.7% vs
60.5%, Table 30) while Alloy 13 demonstrated highest loss in ductility in
punched specimens
10 vs EDM cut (5.2% vs 39.1%, Table 30). High edge damage occurs in punched
specimens
from alloy with lower austenite stability.
Table 8 Tensile Properties of Punched vs EDM Cut Specimens from Selected
Alloys
Allo Cutting Yield Strength
Ultimate Tensile Tensile
y
Method (MPa) Strength (MPa)
Elongation (%)
392 1310 46.7
EDM Cut 397 1318 45.1
400 1304 49.7
Alloy 1
431 699 9.3
Punched 430 680 8.1
422 656 6.9
434 1213 46.4
EDM Cut 452 1207 46.8
444 1199 49.1
Alloy 2
491 823 14.4
Punched 518 792 11.3
508 796 11.9
468 1166 56.1
EDM Cut 480 1177 52.4
Alloy 9 475 1169 56.9
508 1018 29.2
Punched
507 1007 28.6
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Alloy Cutting Yield Strength Ultimate
Tensile Tensile
Method (MPa) Strength (MPa)
Elongation (%)
490 945 23.3
474 1115 64.4
EDM Cut 464 1165 62.5
495 1127 62.7
Alloy 11
503 924 24.6
Punched 508 964 28.0
490 921 25.7
481 1094 54.4
EDM Cut 479 1128 64.7
495 1126 62.4
Alloy 12
521 954 27.1
Punched 468 978 30.7
506 975 31.2
454 1444 39.5
EDM Cut
450 1455 38.7
Alloy 13 486 620 5.0
Punched 469 599 6.3
483 616 4.5
484 1170 58.7
EDM Cut 489 1182 61.2
468 1188 59.0
Alloy 14
536 846 17.0
Punched 480 816 18.4
563 870 17.5
445 1505 37.8
EDM Cut
422 1494 37.5
Alloy 18 478 579 2.4
Punched 469 561 2.6
463 582 2.9
464 1210 57.6
Alloy 21 EDM Cut 499 1244 49.0
516 1220 54.5
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Alloy Cutting Yield Strength Ultimate
Tensile Tensile
Method (MPa) Strength (MPa)
Elongation (%)
527 801 11.3
Punched 511 806 12.6
545 860 15.2
440 1166 31.0
EDM Cut 443 1167 32.0
455 1176 31.0
Alloy 24
496 696 5.0
Punched 463 688 5.0
440 684 4.0
474 1183 15.8
EDM Cut 470 1204 17.0
485 1223 17.4
Alloy 25
503 589 2.1
Punched 517 579 0.8
497 583 2.1
735 1133 20.8
EDM Cut
742 1109 19.0
Alloy 26 722 898 3.4
Punched 747 894 2.9
764 894 3.1
537 1329 19.3
EDM Cut 513 1323 21.4
480 1341 20.8
Alloy 27
563 624 4.3
Punched 568 614 3.3
539 637 4.3
460 1209 54.7
EDM Cut 441 1199 54.1
475 1216 52.9
Alloy 34
489 828 15.4
Punched 486 811 14.6
499 813 14.8
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Alloy Cutting Yield Strength
Ultimate Tensile Tensile
Method (MPa) Strength (MPa)
Elongation (%)
431 1196 50.6
EDM Cut 437 1186 52.0
420 1172 54.7
Alloy 35
471 826 19.9
Punched 452 828 19.7
482 854 19.7
Table 9 Tensile Elongation in Specimens with Different Cutting Methods
Loss In Tensile
Average Tensile Elongation (%)
Elongation
Alloy (E2/E1)
EDM Cut (El) Punched (E2) MM Max
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 from Table 30, EDM cutting is considered to
be representative of the optimal
mechanical properties of the identified alloys, without a sheared edge, and
which were
processed to the point of assuming Structure #4 (Recrystallized Modal
Structure).
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Accordingly, samples having a sheared edge due to punching indicate a
significant drop in
ductility as reflected by tensile elongation measurements of the punched
samples having the
ASTM E8 geometry. For Alloy 1, tensile elongation is initially 47.2% and then
drops to
8.1%, a drop itself of 82.8%%. The drop in ductility from the punched to the
EDM cut
(E2/E1) varies from 0.57 to 0.05.
The edge status after punching and EDM cutting was analyzed by SEM using an
EVO-MA10
scanning electron microscope manufactured by Carl Zeiss SMT Inc. The typical
appearance
of the specimen edge after EDM cutting is shown for Alloy 1 in FIG. 36a. The
EDM cutting
method minimizes the damage of a cut edge allowing the tensile properties of
the material to
be measured without any deleterious edge effects. In wire-EDM cutting,
material is removed
from the edge by a series of rapidly recurring current discharges / sparks and
by this route an
edge is formed without substantial deformation or edge damage. The appearance
of the
sheared edge after punching is shown in FIG. 36b. A significant damage of the
edge occurs
in a fracture zone that undergoes severe deformation during punching leading
to structural
transformation in the shear affected zone into a Refined High Strength
Nanomodal Structure
(FIG. 37b) with limited ductility while Recrystallized Modal Structure was
observed near
EDM cut edge (FIG. 37a).
This Case Example demonstrates that in a case of wire-EDM cutting tensile
properties are
measured at relative higher level as compared to that after punching. In
contrast to EDM
cutting, punching of the tensile specimens creates a significant edge damage
which results in
tensile property decrease. Relative excessive plastic deformation of the sheet
alloys herein
during punching leads to structural transformation to a Refined High Strength
Nanomodal
Structure (Structure #5, FIG. 1B) with reduced ductility leading to premature
cracking at the
edge and relatively lower properties (e.g. reduction in elongation and tensile
strength). The
magnitude of this drop in tensile properties has also been observed to depend
on the alloy
chemistry in correlation with austenite stability.
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Case Example #7 Punched Edge vs EDM Cut Tensile Properties and Recovery
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 31
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 min as described herein. Resultant
sheet from
5 each alloy with final thickness of 1.2 mm and Recrystallized Modal
Structure (Structure #4,
FIG. 1B) was used to demonstrate edge damage recovery by annealing of punched
tensile
specimens. In the broad context of the present invention, annealing may be
achieved by
various methods, including but not limited to furnace heat treatment,
induction heat treatment
and/or laser heat treatment.
10 Tensile specimens in the ASTM E8 geometry were prepared using both wire
EDM cutting
and punching. Part of punched tensile specimens was then put through a
recovery anneal of
850 C for 10 minutes, followed by an air cool, to confirm the ability to
recover properties lost
by punching and shearing damage. Tensile properties were measured on an
Instron 5984
mechanical testing frame using Instron's Bluehill control software. All tests
were conducted
15 at room temperature, with the bottom grip fixed and the top grip set to
travel upwards at a
rate of 0.012 mm/s. Strain data was collected using Instron's Advanced Video
Extensometer.
Tensile testing results are provided in Table 31 and illustrated in FIG. 38
for selected alloys
showing a substantial mechanical property recovery in punched samples after
annealing.
For example, in the case of Alloy 1 indicated, when EDM cut into a tensile
testing sample, a
20 tensile elongation average value is about 47.2%. As noted above, when
punched and
therefore containing a sheared edge, the tensile testing of the sample with
such edge indicated
a significant drop in such elongation values, i.e. an average value of only
about 8.1% due to
Mechanism #4 and formation of Refined High Strength Nanomodal Structure
(Structure #5,
FIG. 1B), which while present largely at the edge section where shearing
occurred, is
25 nonetheless reflected in the bulk property measurements in tensile
testing. However, upon
annealing, which is representative of Mechanism #3 in FIG. 1B and conversion
to Structure
#4 (Recrystallized Modal Structure, FIG. 1B), the tensile elongation
properties are restored.
In the case of Alloy 1, the tensile elongation are brought back to an average
value of about
46.2%. Example tensile stress-strain curves for punched specimens from Alloy 1
with and
30 without annealing are shown in FIG. 39. In Table 32, a summary of the
average tensile
properties and the average lost and gained in tensile elongation is provided.
Note that the
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individual losses and gains are a larger spread than the average losses.
Accordingly, in the
context of the present disclosure, the alloys herein, having an initial value
of tensile
elongation (El) when sheared, may indicate a drop in elongation properties to
a value of E2,
wherein E2= (0Ø57 to 0.05)(E1). Then, upon application of Mechanism #3,
which is
preferably accomplished by heating/annealing at a temperature range of 450 C
up to the Tm
depending on alloy chemistry, the value of E2 is recovered to an elongation
value E3 =(0.48 to
1.21)(E1).
Table 10 Tensile Properties of Punched and Annealed Specimens from Selected
Alloys
Yield Strength Ultimate Tensile Tensile
Alloy Cutting Method
(MPa) Strength (MPa)
Elongation (%)
392 1310 46.7
EDM Cut 397 1318 45.1
400 1304 49.7
431 699 9.3
Alloy 1 Punched 430 680 8.1
422 656 6.9
364 1305 43.6
Punched &
364 1315 47.6
Annealed
370 1305 47.3
434 1213 46.4
EDM Cut 452 1207 46.8
444 1199 49.1
491 823 14.4
Alloy 2 Punched 518 792 11.3
508 796 11.9
432 1205 50.4
Punched &
426 1191 50.7
Annealed
438 1188 49.3
468 1166 56.1
Alloy 9 EDM Cut 480 1177 52.4
475 1169 56.9
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Yield Strength Ultimate Tensile Tensile
Alloy Cutting Method
(MPa) Strength (MPa)
Elongation (%)
508 1018 29.2
Punched 507 1007 28.6
490 945 23.3
411 1166 59.0
Punched &
409 1174 52.7
Annealed
418 1181 55.6
474 1115 64.4
EDM Cut 464 1165 62.5
495 1127 62.7
503 924 24.6
Alloy 11 Punched 508 964 28.0
490 921 25.7
425 1128 64.5
Punched &
429 1117 57.1
Annealed
423 1140 54.3
481 1094 54.4
EDM Cut 479 1128 64.7
495 1126 62.4
521 954 27.1
Alloy 12 Punched 468 978 30.7
506 975 31.2
419 1086 65.7
Punched &
423 1085 63.0
Annealed
415 1100 53.8
454 1444 39.5
EDM Cut
450 1455 38.7
486 620 5.0
Punched 469 599 6.3
Alloy 13
483 616 4.5
397 1432 41.4
Punched &
397 1437 37.4
Annealed
404 1439 40.3
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Yield Strength Ultimate Tensile Tensile
Alloy Cutting Method
(MPa) Strength (MPa)
Elongation (%)
484 1170 58.7
EDM Cut 489 1182 61.2
468 1188 59.0
536 846 17.0
Alloy 14 Punched 480 816 18.4
563 870 17.5
423 1163 58.3
Punched &
412 1168 55.9
Annealed
415 1177 51.5
445 1505 37.8
EDM Cut
422 1494 37.5
478 579 2.4
Punched 469 561 2.6
Alloy 18
463 582 2.9
398 1506 36.3
Punched &
400 1502 40.3
Annealed
392 1518 35.4
464 1210 57.6
EDM Cut 499 1244 49.0
516 1220 54.5
527 801 11.3
Alloy 21 Punched 511 806 12.6
545 860 15.2
409 1195 47.7
Punched &
418 1214 53.8
Annealed
403 1194 51.8
440 1166 31.0
EDM Cut 443 1167 32.0
455 1176 31.0
Alloy 24
496 696 5.0
Punched 463 688 5.0
440 684 4.0
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Yield Strength Ultimate Tensile Tensile
Alloy Cutting Method
(MPa) Strength (MPa)
Elongation (%)
559 1100 22.3
Punched &
581 1113 22.0
Annealed
561 1100 22.3
474 1183 15.8
EDM Cut 470 1204 17.0
485 1223 17.4
503 589 2.1
Alloy 25 Punched 517 579 0.8
497 583 2.1
457 1143 15.4
Punched &
477 1159 14.6
Annealed
423 1178 16.3
735 1133 20.8
EDM Cut
742 1109 19.0
722 898 3.4
Punched 747 894 2.9
Alloy 26
764 894 3.1
715 1112 18.8
Punched &
713 1098 17.8
Annealed
709 931 10.0
537 1329 19.3
EDM Cut 513 1323 21.4
480 1341 20.8
563 624 4.3
Alloy 27 Punched 568 614 3.3
539 637 4.3
505 1324 19.7
Punched &
514 1325 20.0
Annealed
539 1325 19.4
460 1209 54.7
Alloy 29 EDM Cut 441 1199 54.1
475 1216 52.9
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Yield Strength Ultimate Tensile Tensile
Alloy Cutting Method
(MPa) Strength (MPa)
Elongation (%)
489 828 15.4
Punched 486 811 14.6
499 813 14.8
410 1204 53.9
Punched &
410 1220 53.2
Annealed
408 1214 52.3
431 1196 50.6
EDM Cut 437 1186 52.0
420 1172 54.7
471 826 19.9
Alloy 32 Punched 452 828 19.7
482 854 19.7
406 1169 58.1
Punched &
403 1170 51.4
Annealed
405 1176 57.7
Table 32 Summary of Tensile Properties; Loss (E2/E1) and Gain (E3/E1)
Loss In Tensile Elongation Gain in Tensile Elongation
Alloy (E2/E1) (E3/E1)
Min Max Min Max
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
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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 tensile specimens results in edge damage and lowering the tensile
properties of
the material. Plastic deformation of the sheet alloys herein during punching
leads to
structural transformation to a Refined High Strength Nanomodal Structure
(Structure #5,
FIG. 1B) with reduced ductility leading to premature cracking at the edge and
relatively
lower properties (e.g. reduction in elongation and tensile strength). This
Case Example
demonstrates that due to the unique structural reversibility, the edge damage
in the alloys
listed in Table 2 is substantially recoverable by annealing leading back to
Recrystallized
Modal Structure (Structure #4, FIG. 1B) formation with full or partial
property restoration
that depends on alloy chemistry and processing. For example, as exemplified by
Alloy 1,
punching and shearing and creating a sheared edge is observed to reduce
tensile strength from
an average of about 1310 MPa (an EDM cut sample without a sheared/damaged
edge) to an
average value of 678 MPa, a drop of between 45 to 50%. Upon annealing, tensile
strength
recovers to an average value of about 1308 MPa, which is in the range of
greater than or
equal to 95% of the original value of 1310 MPa. Similarly, tensile elongation
is initially at an
average of about 47.1%, dropping to an average value of 8.1%, a decrease of up
to about 80
to 85%, and upon annealing and undergoing what is shown in FIG. 1B as
Mechanism #3,
tensile elongation recovers to an average value of 46.1%, a recovery of
greater than or equal
to 90% of the value of the elongation value of 47.1%.
Case Example #8 Temperature Effect on Recovery and Recrystallization
Slabs with thickness of 50 mm were laboratory cast from Alloy 1 and laboratory
processed
by hot rolling down to thickness of 2 mm and cold rolling with reduction of
approximately
40%. Tensile specimens in the ASTM E8 geometry were prepared by wire EDM cut
from
cold rolled sheet. Part of tensile specimens was annealed for 10 minutes at
different
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temperatures in a range from 450 to 850 C, followed by an air cool. Tensile
properties were
measured on an Instron 5984 mechanical testing frame using Instron's Bluehill
control
software. All tests were conducted at room temperature, with the bottom grip
fixed and the
top grip set to travel upwards at a rate of 0.012 mm/s. Strain data was
collected using
Instron's Advanced Video Extensometer. Tensile testing results are shown in
FIG. 40
demonstrating a transition in deformation behavior depending on annealing
temperature.
During the process of cold rolling, the Dynamic Nanophase Strengthening
(Mechanism #2,
FIG. 1A) or the Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B)
occurs
which involves, once the yield strength is exceeded with increasing strain,
the continuous
transformation of austenite to ferrite plus one or more types of nanoscale
hexagonal phases.
Concurrent with this transformation, deformation by dislocation mechanisms
also occurs in
the matrix grains prior to and after transformation. The result is the change
in the
microstructure from the Nanomodal Structure (Structure #2, FIG. 1A) to the
High Strength
Nanomodal Structure (Structure #3, FIG. 1A) or from the Recrystallized Modal
Structure
(Structure #4, FIG. 1B) to the Refined High Strength Nanomodal Structure
(Structure #5,
FIG. 1B). The structure and property changes occurring during cold deformation
can be
reversed at various degrees by annealing depending on annealing parameters as
seen in the
tensile curves of FIG. 40A. In FIG. 40B, the corresponding yield strength from
the tensile
curves are provided as a function of the heat treatment temperature. The yield
strength after
cold rolling with no anneal is measured at 1141 MPa. As shown, depending on
how the
material is annealed which may include partial and full recovery and partial
and full
recrystallization the yield strength can be varied widely from 1372 MPa at the
500 C anneal
down to 458 MPa at the 850 C anneal.
To show the microstructural recovery in accordance to the tensile property
upon annealing,
TEM studies were conducted on selected samples that were annealed at different
temperatures. For comparison, cold rolled sheet was included as a baseline
herein.
Laboratory cast Alloy 1 slab of 50 mm thick was used, and the slab was hot
rolled at 1250 C
by two-step of 80.8% and 78.3% to approx. 2 mm thick, then cold rolled by 37%
to sheet of
1.2 mm thick. The cold rolled sheet was annealed at 450 C, 600 C, 650 C and
700 C
respectively for 10 minutes. FIG. 41 shows the microstructure of as-cold
rolled Alloy 1
sample. It can be seen that typical High Strength Nanomodal Structure is
formed after cold
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rolling, in which high density of dislocations are generated along with the
presence of strong
texture. Annealing at 450 C for 10 mm does not lead to recrystallization and
formation of
the High Strength Nanomodal Structure, as the microstructure remains similar
to that of the
cold rolled structure and the rolling texture remains unchanged (FIG. 42).
When the cold
rolled sample is annealed at 600 C for 10 min, TEM analysis shows very small
isolated
grains, a sign of the beginning of recrystallization. As shown in FIG. 43,
isolated grains of
100 nm or so are produced after the annealing, while areas of deformed
structure with
dislocation networks are also present. Annealing at 650 C for 10 mm shows
larger
recrystallized grains suggesting the progress of recrystallization. Although
the fraction of
deformed area is reduced, the deformed structure continues to be seen, as
shown in FIG. 44.
Annealing at 700 C 10 mm shows larger and cleaner recrystallized grains, as
displayed by
FIG. 45. Selected electron diffraction shows that these recrystallized grains
are of the
austenite phase. The area of deformed structure is smaller compared to the
samples annealed
at lower temperature. Survey over the entire sample suggests that approx. 10%
to 20% area
is occupied by the deformed structure. The progress of recrystallization
revealed by TEM in
the samples annealed at lower temperature to higher temperature corresponds
excellently to
the change of tensile properties shown in FIG. 40. These low temperature
annealed samples
(such as below 600 C) maintain predominantly the High Strength Nanomodal
Structure,
leading to the reduced ductility. The recrystallized sample (such as at 700 C)
recovers
majority of the elongation, compared to the fully recrystallized sample at 850
C. The
annealing in between these temperatures partially recovers the ductility.
One reason behind the difference in recovery and transition in deformation
behavior is
illustrated by the model TTT diagram in FIG. 46. As described previously, the
very fine /
nanoscale grains of ferrite formed during cold working recrystallize into
austenite during
annealing and some fraction of the nanoprecipitates re-dissolve. Concurrently,
the effect of
the strain hardening is eliminated with dislocation networks and tangles, twin
boundaries, and
small angle boundaries being annihilated by various known mechanisms. As shown
by the
heating curve A of the model temperature, time transformation (TTT) diagram in
FIG. 46, at
low temperatures (particularly below 650 C for Alloy 1), only recovery may
occur without
recrystallization (i.e. recovery being a reference to a reduction in
dislocation density).
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In other words, in the broad context of the present invention, the effect of
shearing and
formation of a sheared edge, and its associated negative influence on
mechanical properties,
can be at least partially recovered at temperatures of 450 C up to 650 C as
shown in FIG. 46.
In addition, at 650 C and up to below Tm of the alloy, recrystallization can
occur, which also
contributes to restoring mechanical strength lost due to the formation of a
sheared edge.
Accordingly, this Case Example demonstrates that upon deformation during cold
rolling,
concurrent processes occur involving dynamic strain hardening and phase
transformation
through unique Mechanisms #2 or #3 (FIG. 1A) along with dislocation based
mechanisms.
Upon heating, the microstructure can be reversed into a Recrystallized Modal
Structure
(Structure #4, FIG. 1B). However, at low temperatures, this reversing process
may not occur
when only dislocation recovery takes place. Thus, due to the unique mechanisms
of the
alloys in Table 2, various external heat treatments can be used to heal the
edge damage from
punching / stamping.
Case Example #9 Temperature Effect of Punched Edge Recovery
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 33
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 min as described in Main Body
section of current
application. Resultant sheet from each alloy with final thickness of 1.2 mm
and
Recrystallized Modal Structure (Structure #4, FIG. 1B) was used to demonstrate
punched
edge damage recovery after annealing as a function of temperature.
Tensile specimens in the ASTM E8 geometry were prepared by punching. A part of
punched
tensile specimens from selected alloys was then put through a recovery anneal
for 10 minutes
at different temperatures in a range from 450 to 850 C, followed by an air
cool. Tensile
properties were measured on an Instron 5984 mechanical testing frame using
Instron' s
Bluehill control software. All tests were conducted at room temperature, with
the bottom
grip fixed and the top grip set to travel upwards at a rate of 0.012 mm/s.
Strain data was
collected using Instron's Advanced Video Extensometer.
Tensile testing results are shown in Table 32 and in FIG. 47. As it can be
seen, full or nearly
full property recovery achieved after annealing at temperatures at 650 C and
higher,
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suggesting that the structure is fully or near fully recrystallized (i.e.
change in structure from
Structure #5 to Structure #4 in FIG. 1B) in the damaged edges after punching.
For example,
the level of recrystallization at the damaged edge is contemplated to be at a
level of greater
than or equal to 90% when annealing temperatures are in the range of 650 C up
to Tm.
5 Lower annealing temperature (e.g. temperatures below 650 C does not
result in full
recrystallization and leads to partial recovery (i.e. decrease in dislocation
density) as
described in Case Example #8 and illustrated in FIG. 6.
Microstructural changes in Alloy 1 at the shear edge as a result of the
punching and annealing
at different temperatures were examined by SEM. Cross section samples were cut
from
10 ASTM E8 punched tensile specimens near the sheared edge in as-punched
condition and after
annealing at 650 C and 700 C as shown in FIG. 48.
For SEM study, the cross section samples were ground on SiC abrasive papers
with reduced
grit size, and then polished progressively with diamond media paste down to 1
um. The final
polishing was done with 0.02 um grit 5i02 solution. Microstructures were
examined by
15 SEM using an EVO-MA10 scanning electron microscope manufactured by Carl
Zeiss SMT
Inc.
FIG. 49 shows the backscattered SEM images of the microstructure at the edge
in the as-
punched condition. It can be seen that the microstructure is deformed and
transformed in the
shear affected zone (i.e., the triangle with white contrast close to the edge)
in contrast to the
20 recrystallized microstructure in the area away from the shear affected
zone. Similar to tensile
deformation, the deformation in the shear affected zone caused by punching
creates Refined
High Strength Nanomodal Structure (Structure #5, FIG. 1B) through Nanophase
Refinement
& Strengthening mechanism. However, annealing recovers the tensile properties
of punched
ASTM E8 specimens, which are related to the microstructure change in the shear
affected
25 zone during annealing. FIG. 50 shows the microstructure of the sample
annealed at 650 C
for 10 minutes. Compared to the as-punched sample, the shear affected zone
becomes
smaller with less contrast suggesting that the microstructure in the shear
affected zone
evolves toward that in the center of the sample. A high magnification SEM
image shows that
some very small grains are nucleated, but recrystallization does not take
place massively
30 across the shear affected zone. It is likely that the recrystallization
is in the early stage with
most of the dislocations annihilated. Although the structure is not fully
recrystallized, the
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tensile property is substantially recovered (Table 32 and FIG.47a). Annealing
at 700 C for
minutes leads to full recrystallization of the shear affected zone. As shown
in FIG. 51, the
contrast in shear affected zone significantly decreased. High magnification
image shows that
equiaxed grains with clear grain boundaries are formed in the shear affected
zone, indicating
5 full recrystallization. The grain size is smaller than that in the center
of sample. Note that the
grains in the center are resulted from recrystallization after annealing at
850 C for 10 minutes
before punching of specimens. With the shear affected zone fully
recrystallized, the tensile
properties are fully recovered, as shown in Table 32 and FIG. 47a.
Punching of tensile specimens result in edge damage lowering the tensile
properties of the
10 material. Plastic deformation of the sheet alloys herein during punching
leads to structural
transformation to a Refined High Strength Nanomodal Structure (Structure #5,
FIG. 1B) with
reduced ductility leading to premature cracking at the edge. This Case Example
demonstrates
that this edge damage is partially / fully recoverable by different anneals
over a wide range of
industrial temperatures.
Table 33 Tensile Properties after Punching and Annealing at Different
Temperatures
Anneal Yield
Ultimate Tensile Tensile Elongation
Alloy Temperature Strength
( C) (MPa) Strength (MPa) (%)
494 798 12.6
As Punched 487 829 14.3
474 792 15.3
481 937 21.5
450 469 934 20.9
485 852 19.3
Alloy 1 464 1055 27.3
600 472 1103 30.5
453 984 23.7
442 1281 51.5
650 454 1270 45.4
445 1264 51.1
700 436 1255 50.1
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Anneal Yield
Ultimate Tensile Tensile Elongation
Alloy Temperature Strength
MP Strength (MPa) (%)
( C) (a)
442 1277 52.1
462 1298 51.6
407 1248 52.0
850 406 1260 47.8
412 1258 48.5
508 1018 29.2
As Punched 507 1007 28.6
490 945 23.3
461 992 28.5
600 462 942 24.8
471 968 25.6
460 1055 33.0
Alloy 9 650 470 1166 48.3
473 1177 49.3
457 1208 57.5
700 455 1169 50.3
454 1171 61.6
411 1166 59.0
850 409 1174 52.7
418 1181 55.6
521 954 27.1
As Punched 468 978 30.7
506 975 31.2
462 1067 44.9
600 446 1013 41.3
Alloy 12 471 1053 41.1
452 1093 61.5
650 449 1126 57.8
505 1123 55.4
480 1112 59.6
700
460 1117 61.8
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Anneal Yield
Ultimate Tensile Tensile Elongation
Alloy Temperature Strength
( C) (MPa) Strength (MPa) (%)
468 1096 61.5
419 1086 65.7
850 423 1085 63.0
415 1100 53.8
Case Example #10 Effect of Punching Speed on Punched Edge Property
Reversibility
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 34
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 min as described herein. Resultant
sheet from
each alloy with final thickness of 1.2 mm and Recrystallized Modal Structure
(Structure #4,
FIG. 1B) was used to demonstrate edge damage recovery as a function of
punching speed.
Tensile specimens in the ASTM E8 geometry were prepared by punching at three
different
speeds of 28 mm/s, 114 mm/s, and 228 mm/s. Wire EDM cut specimens from the
same
materials were used for the reference. A part of punched tensile specimens
from selected
alloys was then put through a recovery anneal for 10 minutes at 850 C,
followed by an air
cool. Tensile properties were measured on an Instron 5984 mechanical testing
frame using
Instron's Bluehill control software. All tests were conducted at room
temperature, with the
bottom grip fixed and the top grip set to travel upwards at a rate of 0.012
mm/s. Strain data
was collected using Instron's Advanced Video Extensometer. Tensile testing
results are
listed in Table 34 and tensile properties as a function of punching speed for
selected alloys
are illustrated in FIG. 52. It is seen that tensile properties drop
significantly in the punched
samples as compared to that for wire EDM cut. Punching speed increase from 28
mm/s to
228 mm/s leads to increase in properties of all three selected alloys. The
localized heat
generation during punching a hole or shearing an edge is known to increase
with increasing
punching velocity and might be a factor in edge damage recovery in specimens
punched at
higher speed. Note that heat alone will not cause edge damage recovery but
will be enabled
by the materials response to the heat generated. This difference in response
for the alloys
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contained in Table 2 in this application to commercial steel samples is
clearly illustrated in
Case Examples 15 and 17.
Table 34 Tensile Properties of Specimens Punched at Different Speed vs EDM Cut
Yield
Sample Preparation Tensile Strength Tensile Elongation
Alloy Strength
Method (MPa) (%)
(MPa)
459 1255 51.2
443 1271 46.4
EDM 441 1248 52.7
453 1251 55.0
467 1259 51.3
474 952 21.8
228 mm/s Punched 498 941 21.6
Alloy 1
493 956 21.6
494 798 13.4
114 mm/s Punched 487 829 15.1
474 792 14.1
464 770 12.8
28 mm/s Punched 479 797 13.7
465 755 12.1
468 1166 56.1
EDM 480 1177 52.4
475 1169 56.9
500 1067 35.1
228 mm/s Punched 493 999 28.8
470 1042 31.8
Alloy 9
508 1018 29.2
114 mm/s Punched 507 1007 28.6
490 945 23.3
473 851 19.7
28 mm/s Punched 472 841 16.4
494 846 18.9
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Sample Preparation YieldTensile Strength Tensile Elongation
Alloy Strength
Method (MPa) (%)
(MPa)
481 1094 54.4
EDM 479 1128 64.7
495 1126 62.4
495 1124 53.8
228 mm/s Punched
484 1123 53.0
Alloy 12 521 954 27.1
114 mm/s Punched 468 978 30.7
506 975 31.2
488 912 23.6
28 mm/s Punched 472 900 21.7
507 928 22.9
This Case Example demonstrates that punching speed can have a significant
effect on the
resulting tensile properties in steel alloys herein. Localized heat generation
at punching
might be a factor in recovery of the structure near the edge leading to
property improvement.
5
Case Example #11 Edge Structure Transformation During Hole Punching and
Expansion
Slabs with thickness of 50 mm were laboratory cast from Alloy 1 and laboratory
processed
by hot rolling, cold rolling and annealing at 850 C for 10 min as described
herein. Resultant
10 sheet with final thickness of 1.2 mm and Recrystallized Modal Structure
(Structure #4, FIG.
1B) was used for hole expansion ratio (HER) tests.
Specimens for testing with a size of 89 x 89 mm were wire EDM cut from the
sheet. The
hole with 10 mm diameter was cut in the middle of specimens by utilizing two
methods:
punching and drilling with edge milling. The hole punching was done on an
Instron Model
15 5985 Universal Testing System using a fixed speed of 0.25 mm/s with 16%
clearance. Hole
expansion ratio (HER) testing was performed on the SP-225 hydraulic press and
consisted of
slowly raising the conical punch that uniformly expanded the hole radially
outward. A digital
image camera system was focused on the conical punch and the edge of the hole
was
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monitored for evidence of crack formation and propagation. The initial
diameter of the hole
was measured twice with calipers, measurements were taken at 90 increments
and averaged
to get the initial hole diameter. The conical punch was raised continuously
until a crack was
observed propagating through the specimen thickness. At that point the test
was stopped and
the hole expansion ratio was calculated as a percentage of the initial hole
diameter measured
before the start of the test. After expansion four diameter measurements were
taken using
calipers every 450 and averaged to account for any asymmetry of the hole due
to cracking.
Results of HER testing are shown in FIG. 53 demonstrating a significantly
lower value for
the sample when the hole was prepared by punching as compared to milling: 5.1%
HER vs
73.6% HER, respectively. Samples were cut from both tested samples as shown in
FIG. 54
for SEM analysis and microhardness measurements.
Microhardness was measured for Alloy 1 at all relevant stages of the hole
expansion process.
Microhardness measurements were taken along cross sections of sheet samples in
the
annealed (before punching and HER testing), as-punched, and HER tested
conditions.
Microhardness was also measured in cold rolled sheet from Alloy 1 for
reference.
Measurement profiles started at an 80 micron distance from the edge of the
sample, with an
additional measurement taken every 120 microns until 10 such measurements were
taken.
After that point, further measurements were taken every 500 microns, until at
least 5 mm of
total sample length had been measured. A schematic illustration of
microhardness
measurement locations in HER tested samples is shown in FIG. 55. SEM images of
the
punched and HER tested samples after microhardness measurements are shown in
FIG. 56.
As shown in FIG. 57, the punching process creates a transformed zone of
approximately 500
microns immediately adjacent to the punched edge, with the material closest to
the punched
edge either fully or near-fully transformed, as evidenced by the hardness
approaching that
observed in the fully-transformed, 40% cold rolled material immediately next
to the punched
edge. Microhardness profiles for each sample is presented in FIG. 58. As it
can be seen,
microhardness gradually increases towards a hole edge in the case of milled
while in the case
of punched hole microhardness increase was observed in a very narrow area
close to the hole
edge. TEM samples were cut at the same distance in both cases as indicated in
FIG. 58.
To prepare the TEM specimens, the HER test samples were first sectioned by
wire EDM, and
a piece with a portion of hole edge was thinned by grinding with pads of
reduced grit size.
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Further thinning to ¨60 um thickness is done by polishing with 9 um, 3 um, and
1 um
diamond suspension solution respectively. Discs of 3 mm in diameter were
punched from the
foils near the edge of the hole and the final polishing was completed by
electropolishing
using a twin-jet polisher. The chemical solution used was a 30% Nitric acid
mixed in
Methanol base. In case of insufficient thin area for TEM observation, the TEM
specimens
may be ion-milled using a Gatan Precision Ion Polishing System (PIPS). The ion-
milling
usually is done at 4.5 keV, and the inclination angle is reduced from 4 to 2
to open up the
thin area. The TEM studies were done using a JEOL 2100 high-resolution
microscope
operated at 200 kV. Since the location for TEM study is at the center of the
disc, the
observed microstructure is approximately ¨1.5 mm from the edge of hole.
The initial microstructure of the Alloy 1 sheet before testing is shown on
FIG. 59
representing Recrystallized Modal Structure (Structure #4, FIG. 1B). FIG. 60a
shows the
TEM micrograph of the microstructure in the HER test sample with punched hole
after
testing (HER = 5.1%) in different areas at the location of 1.5 mm from hole
edge. It was
found that mainly the recrystallized microstructure remains in the sample
(FIG. 60a) with
small amount of area with partially transformed "pockets" (FIG. 60b)
indicating that limited
volume (¨ 1500 um deep) of the sample was involved in deformation at HER
testing. In the
HER sample with milled hole (HER = 73.6%), as shown in FIG. 61, there is a
great amount
of deformation in the sample as indicated by a large amount of transformed
"pockets" and
high density of dislocations (108 to 1010 mm-2).
To analyze in more detail the reason causing the poor HER performance in
samples with
punched holes, Focused Ion Beam (FIB) technique was utilized to make TEM
specimens at
the very edge of the punched hole. As shown in FIG. 62, TEM specimen is cut at
¨10 um
from the edge. To prepare TEM specimens by FIB, a thin layer of platinum is
deposited on
the area to protect the specimen to be cut. A wedge specimen is then cut out
and lifted by a
tungsten needle. Further ion milling is performed to thin the specimen.
Finally the thinned
specimen is transferred and welded to copper grid for TEM observation. FIG. 63
shows the
microstructure of the Alloy 1 sheet at the distance of ¨10 micron from the
punched hole edge
which is significantly refined and transformed as compared to the
microstructure in the Alloy
1 sheet before punching. It suggests that punching caused severe deformation
at the hole
edge such that Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B)
occurred
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leading to formation of Refined High Strength Nanomodal Structure (Structure
#5, FIG. 1B)
in the area close to the punched hole edge. This structure has relative lower
ductility as
compared to Recrystallized Modal Structure Table 1 resulting in premature
cracking at the
edge and low HER values. This Case Example demonstrates that the alloys in
Table 2
exhibit the unique ability to transform from a Recrystallized Modal Structure
(Structure #4,
FIG. 1B) to a Refined High Strength Nanomodal Structure (Structure #5, FIG.
1B) through
the identified Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B).
The
structural transformation occurring due to deformation at the hole edge at
punching appears
to be similar in nature to transformation occurring during cold rolling
deformation and that
observed during tensile testing deformation.
Case Example #12 HER Testing Results With and Without Annealing
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 35
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 min as described herein. Resultant
sheet with final
thickness of 1.2 mm and Recrystallized Modal Structure (Structure #4, FIG. 1B)
was used for
hole expansion ratio (HER) tests.
Test specimens of 89 x 89 mm were wire EDM cut from the sheet from larger
sections. A 10
mm diameter hole was made in the center of specimens by punching on an Instron
Model
5985 Universal Testing System using a fixed speed of 0.25 mm/s at 16% punch to
die
clearance. Half of the prepared specimens with punched holes were individually
wrapped in
stainless steel foil and annealed at 850 C for 10 minutes before HER testing.
Hole expansion
ratio (HER) testing was performed on the SP-225 hydraulic press and consisted
of slowly
raising the conical punch that uniformly expanded the hole radially outward. A
digital image
camera system was focused on the conical punch and the edge of the hole was
monitored for
evidence of crack formation and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
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measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
. The results of the hole expansion ratio measurements on the specimens with
and without
annealing after hole punching are shown in Table 35. As shown in FIG. 64, FIG.
65, FIG.
66, FIG. 67 and FIG. 68 for Alloy 1, Alloy 9, Alloy 12, Alloy 13, and Alloy
17, respectively,
the hole expansion ratio measured with punched holes with annealing is
generally greater
than in punched holes without annealing. The increase in hole expansion ratio
with annealing
for the identified alloys herein therefore leads to an increase in the actual
HER of about 25%
to 90%.
Table 35 Hole Expansion Ratio Results for Select Alloys With and Without
Annealing
Punch Measured Hole Average Hole
Clearance Expansion Ratio Expansion Ratio
Material Condition (%) (%) (%)
3.00
Without
16 3.90 3.20
Annealing
2.70
Alloy 1
105.89
With
16 81.32 93.10
Annealing
92.11
3.09
Without
16 3.19 3.19
Annealing
3.29
Alloy 9
78.52
With
16 97.60 87.84
Annealing
87.40
Without 4.61
16 4.91
Annealing 5.21
Alloy 12
With 69.11
16 77.60
Annealing 83.60
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Clearance Expansion Ratio Expansion Ratio
Material Condition (%) (%) (%)
80.08
1.70
Without
16 1.40 1.53
Annealing
1.50
Alloy 13
32.37
With
16 29.00 31.12
Annealing
32.00
12.89
Without
16 28.70 21.46
Annealing
22.80
Alloy 17
104.21
With
16 80.42 103.74
Annealing
126.58
This Case Example demonstrates that edge formability demonstrated during HER
testing can
yield poor results due to edge damage during the punching operation as a
result of the unique
mechanisms in the alloys listed in Table 2. The fully post processed alloys
exhibit very high
5 tensile ductility as shown in Table 6 through Table 10 coupled with very
high strain
hardening and resistance to necking until near failure. Thus, the material
resists catastrophic
failure to a great extent but during punching, artificial catastrophic failure
is forced to occur
near the punched edge. Due to the unique reversibility of the identified
mechanisms, this
deleterious edge damage as a result of Nanophase Refinement & Strengthening
(Mechanism
10 #3, FIG. 1A) and structural transformation can be reversed by annealing
resulting in high
HER results. Thus, high hole expansion ratio values can be obtained in a case
of punching
hole with following annealing and retaining exceptional combinations of
tensile properties
and the associated bulk formability.
In addition, it can be appreciated that the alloys herein that have undergone
the processing
15 pathways to provide such alloys in the form of Structure #4
(Recrystallized Modal Structure)
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will indicate, for a hole that is formed by shearing, and including a sheared
edge, a first hole
expansion ratio (HERD and upon heating the alloy will have a second hole
expansion ratio
(HER2), wherein HER2>HER1.
More specifically, it can also be appreciated that the alloys herein that have
undergone the
processing pathways to provide such alloys with Structure #4 (Recrystallized
Modal
Structure) will indicate, for a hole that was placed in the alloy through
methods (i.e. waterjet
cutting, laser cutting, wire-edm, milling etc.) where the hole that is formed
that does not rely
primarily on shearing, compared to punching a hole, a first hole expansion
ratio (HERO
where such value may itself fall in the range of 30 to 130%. However, when the
same alloy
includes a hole formed by shearing, a second hole expansion ratio is observed
(HER2)
wherein HER2 = (0.01 to 0.30)(HER1). However, if the alloy is then subject to
heat treatment
herein, it is observed that HER2 is recovered to a HER3= (0.60 to 1.0) HER1.
Case Example #13 Edge Condition Effect on Alloy Properties
Slabs with thickness of 50 mm were laboratory cast from Alloy 1 according to
the atomic
ratios provided in Table 2 and laboratory processed by hot rolling, cold
rolling and annealing
at 850 C for 10 min as described herein. Resultant sheet from Alloy 1 with
final thickness of
1.2 mm and Recrystallized Modal Structure (Structure #4, FIG. 1B) was used to
demonstrate
the effect that edge condition has on Alloy 1 tensile and hole expansion
properties.
Tensile specimens of ASTM E8 geometry were created using two methods: Punching
and
wire EDM cutting. Punched tensile specimens were created using a commercial
press. A
subset of punched tensile specimens was heat treated at 850 C for 10 minutes
to create
samples with a punched then annealed edge condition.
Tensile properties of ASTM E8 specimens were measured on an Instron 5984
mechanical
testing frame using Instron' s Bluehill 3 control software. All tests were
conducted at room
temperature, with the bottom grip fixed and the top grip set to travel upwards
at a rate of
0.025 mm/s for the first 0.5% elongation, and at a rate of 0.125 mm/s after
that point. Strain
data was collected using Instron's Advanced Video Extensometer. Tensile
properties of
Alloy 1 with punched, EDM cut, and punched then annealed edge conditions are
shown in
Table 36. Tensile properties of Alloy 1 with different edge conditions are
shown in FIG. 69.
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Table 36 Tensile Properties of Alloy 1 with Different Edge Conditions
Ultimate Tensile
Edge Tensile
Strength
Condition Elongation (%)
(MPa)
12.6 798
Punched 14.3 829
15.3 792
50.5 1252
51.2 1255
52.7 1248
EDM Cut
55.0 1251
51.3 1259
50.5 1265
52.0 1248
Punched
Then 47.8 1260
Annealed
48.5 1258
Specimens for hole expansion ratio testing with a size of 89 x 89 mm were wire
EDM cut
from the sheet. The holes with 10 mm diameter were prepared by two methods:
punching
and cutting by wire EDM. The punched holes with 10 mm diameter were created by
punching at 0.25 mm/s on an Instron 5985 Universal Testing System with a 16%
punch
clearance and with using the flat punch profile geometry. A subset of punched
samples for
hole expansion testing were annealed with an 850 C for 10 minutes heat
treatment after
punching.
Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
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The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
Hole expansion ratio testing results are shown in Table 37. An average hole
expansion ratio
value for each edge condition is also shown. The average hole expansion ratio
for each edge
condition is plotted in FIG. 70. It can be seen that for samples with EDM cut
and punched
then annealed edge conditions the edge formability (i.e. HER response) is
excellent, whereas
samples with holes in the punched edge condition have considerably lower edge
formability.
Table 37 Hole Expansion Ratio of Alloy 1 with Different Edge Conditions
Measured Hole Average Hole
Edge
Expansion Ratio Expansion Ratio
Condition
(%) (%)
3.00
Punched 3.90 3.20
2.70
92.88
EDM Cut 67.94 82.43
86.47
105.90
Punched
Then 81.30 93.10
Annealed
92.10
This Case Example demonstrates that the edge condition of Alloy 1 has a
distinct effect on
the tensile properties and edge formability (i.e. HER response). Tensile
samples tested with
punched edge condition have diminished properties when compared to both wire
EDM cut
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and punched after subsequent annealing. Samples having the punched edge
condition have
hole expansion ratios averaging 3.20%, whereas EDM cut and punched then
annealed edge
conditions have hole expansion ratios of 82.43% and 93.10%, respectively.
Comparison of
edge conditions also demonstrates that damage associated with edge creation
(i.e. via
punching) has a non-trivial effect on the edge formability of the alloys
herein.
Case Example #14 HER Results as a Function of Hole Punching Speed
Slabs with thickness of 50 mm were laboratory cast from selected alloys listed
in Table 38
according to the atomic ratios provided in Table 2 and laboratory processed by
hot rolling,
cold rolling and annealing at 850 C for 10 min as described herein. Resultant
sheet from
each alloy with final thickness of 1.2 mm and Recrystallized Modal Structure
(Structure #4,
FIG. 1B) were used to demonstrate an effect of hole punching speed on HER
results.
Specimens for testing with a size of 89 x 89 mm were wire EDM cut from the
sheet. The
holes with 10 mm diameter were punched at different speeds on two different
machines but
all of the specimens were punched with a 16% punch clearance and with the same
punch
profile geometry. The low speed punched holes (0.25 mm/s, 8 mm/s) were punched
using an
Instron 5985 Universal Testing System and the high speed punched holes (28
mm/s, 114
mm/s, 228 mm/s) were punched on a commercial punch press. All holes were
punched using
a flat punch geometry.
Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
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Hole expansion ratio values for tests are shown in Table 37. An average hole
expansion
value is shown for each speed and alloy tested at 16% punch clearance. The
average hole
expansion ratio as a function of punch speed is shown in FIG. 71, FIG. 72 and
FIG. 73 for
Alloy 1, Alloy 9, and Alloy 12, respectively. It can be seen that as punch
speed increases, all
alloys tested had a positive edge formability response, as demonstrated by an
increase in hole
expansion ratio. The reason for this increase is believed to be related to the
following effects.
With higher punch speed, the amount of heat generated at the sheared edge is
expected to
increase and the localized temperature spike may result in an annealing effect
(i.e. in-situ
annealing). Alternatively, with increasing punch speed, there may be a reduced
amount of
material transforming from the Recrystallized Modal Structure (i.e. Structure
#4 in Fig. 1B)
to the Refined High Strength Nanomodal Structure (i.e. Structure #5 in Fig.
1B).
Concurrently, the amount of Refined High Strength Nanomodal Structure (i.e.
Structure #5 in
Fig. 1B) may be reduced due to the temperature spike enabling localized
recrystallization (i.e.
Mechanism #3 in Fig. 1B).
Table 38 Hole Expansion Ratio at Different Punch Speeds
Measured Hole Average Hole
Punch Speed
Material Expansion Ratio Expansion Ratio
(mm/s)
(%) (%)
0.25 3.00
0.25 3.90 3.20
0.25 2.70
8 4.49
8 3.49 3.82
Alloy 1 8 3.49
28 8.18
28 8.08 7.74
28 6.97
114 17.03
17.53
114 19.62
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Measured Hole Average Hole
Punch Speed
Material (mm/s) Expansion Ratio Expansion Ratio
(%) (%)
114 15.94
228 20.44
228 21.24 21.70
228 23.41
0.25 3.09
0.25 3.19 3.19
0.25 3.29
8 6.80
8 7.39 6.93
8 6.59
28 21.04
Alloy 9 28 17.35 19.11
28 18.94
114 24.80
114 19.74 24.29
114 28.34
228 26.00
228 35.16 30.57
228 30.55
0.25 4.61
4.91
0.25 5.21
8 7.62
Alloy 12
8 14.61 11.28
8 11.62
28 29.38 31.59
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Measured Hole Average Hole
Punch Speed
Material Expansion Ratio Expansion Ratio
(mm/s)
(%) (%)
28 33.70
28 31.70
114 40.08
114 48.11 45.50
114 48.31
228 50.00
228 40.56 49.36
228 57.51
This Case Example demonstrates a dependence of edge formability on punching
speed as
measured by the hole expansion ratio. As punch speed increases, the hole
expansion ratio
generally increases for the alloys tested. With increased punching speed, the
nature of the
.. edge is changed such that improved edge formability (i.e. HER response) is
achieved. At
punching speeds greater than those measured, edge formability is expected to
continue
improving towards even higher hole expansion ratio values.
Case Example #15 HER in DP980 as a Function of Hole Punching Speed
Commercially produced and processed Dual Phase 980 steel was purchased and
hole
expansion ratio testing was performed. All specimens were tested in the as
received
(commercially processed) condition.
Specimens for testing with a size of 89 x 89 mm were wire EDM cut from the
sheet. The
holes with 10 mm diameter were punched at different speeds on two different
machines but
all of the specimens were punched with a 16% punch clearance and with the same
punch
profile geometry using a commercial punch press. The low speed punched holes
(0.25 mm/s)
were punched using an Instron 5985 Universal Testing System and the high speed
punched
holes (28 mm/s, 114 mm/s, 228 mm/s) were punched on a commercial punch press.
All holes
were punched using a flat punch geometry.
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Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
.. The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
Values for hole expansion tests are shown in Table 39. The average hole
expansion value for
each punching speed is also shown for commercial Dual Phase 980 material at
16% punch
clearance. The average hole expansion value is plotted as a function of
punching speed for
commercial Dual Phase 980 steel in FIG. 74.
Table 39 Hole Expansion Ratio of Dual Phase 980 Steel at Different Punch
Speeds
Measured Hole Average Hole
Punch Speed
Material Expansion Ratio Expansion Ratio
(mm/s)
(%) (%)
0.25 23.55
0.25 20.96 22.45
0.25 22.85
28 18.95
Commercial Dual 17.63
28 18.26
Phase 980
28 18.21
114 17.40
114 23.66 20.09
114 19.22
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Measured Hole Average Hole
Punch Speed
Material Expansion Ratio Expansion Ratio
(mm/s)
(%) (%)
228 27.21
228 24.30 23.83
228 19.98
This Case Example demonstrates that no edge performance effect based on punch
speed is
measureable in Dual Phase 980 steel. For all punch speeds measured on Dual
Phase 980 steel
the edge performance (i.e. HER response) is consistently within the 21% 3%
range,
indicating that edge performance in conventional AHSS is not improved by punch
speed as
expected since the unique structures and mechanisms present in this
application as for
example in Figures la and lb are not present.
Case Example #16: HER Results as a Function of Punch Design
Slabs with thickness of 50 mm were laboratory cast from Alloys 1, 9, and 12
according to the
atomic ratios provided in Table 2 and laboratory processed by hot rolling,
cold rolling and
annealing at 850 C for 10 min as described herein. Resultant sheet from each
alloy with final
thickness of 1.2 mm and Recrystallized Modal Structure (Structure #4, FIG. 1B)
was used to
demonstrate an effect of hole punching speed on HER results.
Tested specimens of 89 x 89 mm were wire EDM cut from larger sections. A 10 mm
diameter hole was punched in the center of the specimen at three different
speeds, 28 mm/s,
114 mm/s, and 228 mm/s at 16% punch clearance and with four punch profile
geometries
using a commercial punch press. These punch geometries used were flat, 6
tapered, 7
conical, and conical flat. Schematic drawings of the 6 tapered, 7 conical,
and conical flat
punch geometries are shown in FIG. 75.
Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
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The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
.. initial hole diameter measured before the start of the test. After
expansion four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
Hole expansion ratio data is included respectively in Table 40, Table 41, and
Table 42 for
Alloy 1, Alloy 9, and Alloy 12 at four punch geometries and at two different
punch speeds.
The average hole expansion values for Alloy 1, Alloy 9, and Alloy 12 are shown
in FIG. 76,
FIG. 77 and FIG. 78, respectively. For all alloys tested, the 7 conical punch
geometry
resulted in the largest or tied for the largest hole expansion ratio compared
to all other punch
geometries. Increased punch speed is also shown to improve the edge
formability (i.e. HER
response) for all punch geometries. At increased punching speed with different
punch
geometries, the alloys herein may be able to undergo some amount of
Recrystallization
(Mechanism #3) as it is contemplated that there could be localized heating at
the edge at such
higher relative punch speeds, triggering Mechanism #3 and formation of some
amount of
Structure #4.
Table 40 Hole Expansion Ratio of Alloy 1 with Different Punch Geometries
Measured Hole Average Hole
Punch Speed
Punch Geometry ( /s Expansion Ratio Expansion Ratio
mm)
(%) (%)
Flat 28 8.18
Flat 28 8.08 7.74
Flat 28 6.97
Flat 114 17.03
Flat 114 19.62 17.53
Flat 114 15.94
Flat 228 20.44 21.70
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Measured Hole Average Hole
Punch Speed
Punch Geometry (mm/) Expansion Ratio Expansion Ratio
s
(%) (%)
Flat 228 21.24
Flat 228 23.41
6 Taper 28 7.87
8.32
6 Taper 28 8.77
6 Taper 114 19.84
6 Taper 114 16.55 18.48
6 Taper 114 19.04
7 Conical 28 8.37
7 Conical 28 12.05 10.56
7 Conical 28 11.25
7 Conical 114 23.41
7 Conical 114 21.14 22.85
7 Conical 114 24.00
7 Conical 228 21.71
7 Conical 228 19.50 21.37
7 Conical 228 22.91
Conical Flat 28 8.47
Conical Flat 28 13.25 11.95
Conical Flat 28 14.14
Conical Flat 114 20.42
Conical Flat 114 19.22 19.75
Conical Flat 114 19.62
Conical Flat 228 24.13
Conical Flat 228 23.31 22.39
Conical Flat 228 19.72
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Table 41 Hole Expansion Ratio of Alloy 9 with Different Punch Geometries
Measured Hole Average Hole
Punch Speed
Punch Geometry (mm/s) Expansion Ratio Expansion Ratio
(%) (%)
Flat 28 21.04
Flat 28 17.35 19.11
Flat 28 18.94
Flat 114 24.80
Flat 114 19.74 24.29
Flat 114 28.34
Flat 228 26.00
Flat 228 35.16 30.57
Flat 228 30.55
6 Taper 28 17.35
6 Taper 28 19.06 19.36
6 Taper 28 21.66
6 Taper 114 29.64
6 Taper 114 32.14 31.14
6 Taper 114 31.64
7 Conical 28 22.63
7 Conical 28 23.61 24.05
7 Conical 28 25.92
7 Conical 114 34.36
7 Conical 114 31.67 32.60
7 Conical 114 31.77
7 Conical 228 36.28
7 Conical 228 38.87 36.44
7 Conical 228 34.16
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Measured Hole Average Hole
Punch Speed
Punch Geometry (mm/) Expansion Ratio Expansion Ratio
s
(%) (%)
Conical Flat 28 27.72
Conical Flat 28 24.63 25.59
Conical Flat 28 24.43
Conical Flat 114 30.28
Conical Flat 114 32.87 32.64
Conical Flat 114 34.76
Conical Flat 228 32.90
Conical Flat 228 37.45 35.45
Conical Flat 228 35.99
Table 42 Hole Expansion Ratio of Alloy 12 with Different Punch Geometries
Measured Hole Average Hole
Punch Speed
Punch Geometry (mm/s) Expansion Ratio Expansion Ratio
Flat 28 29.38
Flat 28 33.70 31.59
Flat 28 31.70
Flat 114 40.08
Flat 114 48.11 45.50
Flat 114 48.31
Flat 228 50.00
Flat 228 40.56 49.36
Flat 228 57.51
6 Taper 28 29.91
30.67
6 Taper 28 32.50
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Measured Hole Average Hole
Punch Speed
Punch Geometry Expansion Ratio Expansion Ratio
(m
m/s)
(%) (%)
6 Taper 28 29.61
6 Taper 114 38.42
6 Taper 114 44.37 41.19
6 Taper 114 40.78
7 Conical 28 34.90
7 Conical 28 33.00 33.76
7 Conical 28 33.37
7 Conical 114 45.72
7 Conical 114 49.30 49.10
7 Conical 114 52.29
7 Conical 228 58.90
7 Conical 228 53.43 54.36
7 Conical 228 50.75
Conical Flat 28 37.15
Conical Flat 28 31.47 34.43
Conical Flat 28 34.66
Conical Flat 114 45.76
Conical Flat 114 45.96 46.36
Conical Flat 114 47.36
Conical Flat 228 57.51
Conical Flat 228 53.48 54.11
Conical Flat 228 51.34
This Case Example demonstrates that for all alloys tested, there is an effect
of punch
geometry on edge formability. For all alloys tested, the conical punch shapes
resulted in the
largest hole expansion ratios, thereby demonstrating that modifying the punch
geometry from
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a flat punch to a conical punch shape reduces the damage within the material
due to the
punched edge and improves edge formability. The 7 conical punch geometry
resulted in the
greatest edge formability increase overall when compared to the flat punch
geometry with the
conical flat geometry producing slightly lower hole expansion ratios across
the majority of
alloys tested. For Alloy 1 the effect of punch geometry is diminished with
increasing
punching speed, with the three tested geometries resulting in nearly equal
edge formability as
measured by hole expansion ratio (FIG. 79). Punch geometry, coupled with
increased punch
speeds have been demonstrated to greatly reduce residual damage from punching
within the
edge of the material, thereby improving edge formability. With higher punch
speed, the
amount of heat generated at the sheared edge is expected to increase and the
localized
temperature spike may result in an annealing effect (i.e. in-situ annealing).
Alternatively,
with increasing punch speed, there may be a reduced amount of material
transforming from
the Recrystallized Modal Structure (i.e. Structure #4 in Fig. 1B) to the
Refined High Strength
Nanomodal Structure (i.e. Structure #5 in Fig. 1B). Concurrently, the amount
of Refined
High Strength Nanomodal Structure (i.e. Structure #5 in Fig. 1B) may be
reduced due to the
temperature spike enabling localized recrystallization (i.e. Mechanism #3 in
Fig. 1B).
Case Example #17: HER in Commercial Steel Grades as a Function of Hole
Punching
Speed
Hole expansion ratio testing was performed on commercial steel grades 780, 980
and 1180.
All specimens were tested in the as received (commercially processed) sheet
condition.
Specimens for testing with a size of 89 x 89 mm were wire EDM cut from the
sheet of each
grade. The holes with 10 mm diameter were punched at different speeds on two
different
machines with the same punch profile geometry using a commercial punch press.
The low
speed punched holes (0.25 mm/s) were punched using an Instron 5985 Universal
Testing
System at 12% clearance and the high speed punched holes (28 mm/s, 114 mm/s,
228 mm/s)
were punched on a commercial punch press at 16% clearance. All holes were
punched using
a flat punch geometry.
Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
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outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
Results from hole expansion tests are shown in Table 43 through Table 45 and
illustrated in
FIG. 80. As it can be seen, the hole expansion ratio does not show improvement
with
increasing punching speed in all tested grades.
Table 43 Hole Expansion Ratio of 780 Steel Grade at Different Punch Speeds
Sample Punch Speed Punch to die Punch
HER
(mm/s) clearance ( %) Geometry
1 12% Flat 44.74
2 5 mm/s 12% Flat 39.42
3 12% Flat 44.57
1 16% Flat 35.22
2 28 mm/s 16% Flat 28.4
3 16% Flat 36.38
1 16% Flat 31.58
2 114 mm/s 16% Flat 33.9
3 16% Flat 22.29
1 16% Flat 31.08
2 228 mm/s 16% Flat 31.85
3 16% Flat 31.31
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Table 44 Hole Expansion Ratio of 980 Steel Grade at Different Punch Speeds
Sample Punch Speed Punch to die Punch
HER
# (mm/s) clearance ( %) Geometry
1 12% Flat 33.73
mm/s
2 12% Flat 35.02
1 16% Flat 26.88
2 28 mm/s 16% Flat 26.44
3 16% Flat 23.83
1 16% Flat 26.81
2 114 mm/s 16% Flat 30.56
3 16% Flat 29.24
1 16% Flat 30.06
2 228 mm/s 16% Flat 30.98
3 16% Flat 30.62
Table 45 Hole Expansion Ratio of 1180 Steel Grade at Different Punch Speeds
Sample Punch Speed Punch to die Punch
HER
# (mm/s) clearance ( %) Geometry
1 12% Flat 26.73
2 5 mm/s 12% Flat 32.9
3 12% Flat 25.4
1 16% Flat 35.32
2 28 mm/s 16% Flat 32.11
3 16% Flat 36.37
1 16% Flat 35.15
2 114 mm/s 16% Flat 30.92
3 16% Flat 32.27
1 228 mm/s 16% Flat 27.25
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Sample Punch Speed Punch to die Punch
HER
(mm/s) clearance (%) Geometry
2 16% Flat 23.98
3 16% Flat 31.18
This Case Example demonstrates that no edge performance effect based on hole
punch speed
is measureable in tested commercial steel grades indicating that edge
performance in
conventional AHSS is not effected or improved by punch speed as expected since
the unique
structures and mechanisms present in this application as for example in FIG.
lA and FIG. 1B
are not present.
Case Example #18: Relationship of post uniform elongation to hole expansion
ratio
Existing steel materials have been shown to exhibit a strong correlation of
the measured hole
expansion ratio and the material's post uniform elongation. The post uniform
elongation of a
material is defined as a difference between the total elongation of a sample
during tensile
testing and the uniform elongation, typically at the ultimate tensile strength
during tensile
testing. Uniaxial tensile testing and hole expansion ratio testing were
completed on Alloy
land Alloy 9 on the sheet material at approximately 1.2 mm thickness for
comparison to
existing material correlations.
Slabs with thickness of 50 mm were laboratory cast of Alloy 1 and Alloy 9
according to the
atomic ratios provided in Table 2 and laboratory processed by hot rolling,
cold rolling
annealing at 850 C for 10 min as described in the Main Body section of this
application.
Tensile specimens in the ASTM E8 geometry were prepared by wire EDM. All
samples
were tested in accordance with the standard testing procedure described in the
Main Body of
this document. An average of the uniform elongation and total elongation for
each alloy were
used to calculate the post uniform elongation. The average uniform elongation,
average total
elongation, and calculated post uniform elongation for Alloy 1 and Alloy 9 are
provided in
Table 46.
Specimens for hole expansion ratio testing with a size of 89 x 89 mm were wire
EDM cut
from the sheet of Alloy 1 and Alloy 9. Holes of 10 mm diameter were punched at
0.25 mm/s
on an Instron 5985 Universal Testing System at 12% clearance. All holes were
punched
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using a flat punch geometry. These test parameters were selected as they are
commonly used
by industry and academic professionals for hole expansion ratio testing.
Hole expansion ratio (HER) testing was performed on the SP-225 hydraulic press
and
consisted of slowly raising the conical punch that uniformly expanded the hole
radially
outward. A digital image camera system was focused on the conical punch and
the edge of
the hole was monitored for evidence of crack formation and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
The measured hole expansion ratio values for Alloy 1 and Alloy 9 are provided
in Table 46.
Table 46 Uniaxial Tensile and Hole Expansion Data for Alloy 1 and Alloy 9
Average Average Post Uniform Hole
Alloy Uniform Total Elongation Expansion
Elongation Elongation (Epul) Ratio
(%) (%) (%) (%)
Alloy 1 47.19 49.29 2.10 2.30
Alloy 9 50.83 56.99 6.16 2.83
Commercial reference data is shown for comparison in Table 47 from [Paul S.K.,
J Mater
Eng Perform 2014; 23:3610.1. For commercial data, S.K. Paul's prediction
states that the
hole expansion ratio of a material is proportional to 7.5 times the post
uniform elongation
(See Equation 1).
HER = 7.5(Epui) Equation 1
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Table 47 Reference Data from [Paul S.K., J Mater Eng Perform 2014;23:3610.[
Post Uniform Hole
Commercial Uniform Total
Elongation Expansion
Steel Grade Elongation Elongation
(Epul) Ratio
(%) (%) (%) (%)
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
TW1P940 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
The Alloy 1 and Alloy 9 post uniform elongation and hole expansion ratio are
plotted in FIG.
81 with the commercial alloy data and S.K. Paul's predicted correlation. Note
that the data
for Alloy 1 and Alloy 9 do not follow the predicted correlation line.
This Case Example demonstrates that for the steel alloys herein, the
correlation between post
uniform elongation and the hole expansion ratio does not follow that for
commercial steel
grades. The measured hole expansion ratio for Alloy 1 and Alloy 9 is much
smaller than the
.. predicted values based on correlation for existing commercial steel grades
indicating an effect
of the unique structures and mechanisms are present in the steel alloys herein
as for example
shown in FIG. la and FIG. lb.
Case Example #19 HER Performance as a Function of Hole Expansion Speed
Slabs with thickness of 50 mm were laboratory cast from three selected alloys
according to
the atomic ratios provided in Table 2 and laboratory processed by hot rolling,
cold rolling and
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annealing at 850 C for 10 min as described herein. Sheet from each alloy
possessing the
Recrystallized Modal Structure with final thickness of 1.2 mm were used to
demonstrate the
effect of hole expansion speed on HER performance.
Specimens for testing with a size of 89 x 89 mm were cut via wire EDM from the
sheet.
Holes of 10 mm diameter were punched at a constant speed of 228 mm/s on a
commercial
punch press. All holes were punched with a flat punch geometry, and with
approximately
16% punch to die clearance.
Hole expansion ratio (HER) testing was performed on an Interlaken Technologies
SP-225
hydraulic press and consisted of raising the conical punch that uniformly
expanded the hole
radially outward. Four hole expansion speeds, synonymous with the conical ram
travel
speed, were used; 5, 25, 50, and 100 mm/min. A digital image camera system was
focused
on the conical punch and the edge of the hole was monitored for evidence of
crack formation
and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
measurements were taken using calipers every 45 and averaged to account for
any
.. asymmetry of the hole due to cracking.
Hole expansion ratio values for the tests are shown in Table 48. The average
hole expansion
ratio value is shown for each speed and alloy tested showing an increase in
HER values with
increasing hole expansion speed in all three alloys. The effect of hole
expansion speed is also
demonstrated in FIG. 82, FIG. 83, and FIG. 84 for Alloy 1, Alloy 9, and Alloy
12,
respectively.
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Table 48 Hole Expansion Ratio in Selected Alloys at Different Expansion Speeds
Hole Expansion Measured Hole Average Hole
Punch Speed
Material (mm/ Speed Expansion Ratio Expansion Ratio
s)
(mm/min) (%) (%)
19.09
22.54 20.55
20.02
30.70
25 29.14 28.58
Alloy 1 228 25.91
34.05
50 36.43 34.63
33.42
37.11
100 38.52 37.19
35.93
34.06
5 34.07 34.15
34.31
32.87
25 45.46 40.77
Alloy 9 228 43.98
38.39
50 39.71 44.17
54.42
48.01
100 55.27 49.50
45.23
48.61
5 34.79 43.51
47.14
42.13
Alloy 12 228 25 57.82 50.64
51.96
63.77
50 68.46 62.97
56.68
100 57.79 56.73
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Hole Expansion Measured Hole Average Hole
Punch Speed
Material (mm/ Speed
Expansion Ratio Expansion Ratio
s)
(mm/min) (%) (%)
49.28
63.11
This Case Example demonstrates that formability of the edge, i.e. its ability
to be deformed
with relatively reduced cracking, as measured by HER testing, can be affected
by the speed
of deformation of the hole edge (i.e. hole expansion speed). The alloys tested
in this Case
Example demonstrated a positive correlation between hole expansion ratio and
the hole
expansion speed, with increasing hole expansion speed resulting in relatively
higher
measured hole expansion ratios.
Accordingly, in the broad context of the present disclosure, it has been
established
that once an edge is formed, of any geometry by any edge formation method
which causes
deformation of the metal alloy when forming the edge (e.g. by punching,
shearing, piercing,
perforating, cutting, cropping, stamping,), by increasing the speed at which
that edge once
formed is then expanded, one observes that the edge itself is then capable of
more expansion
with a relatively reduced tendency to crack. The edge herein can therefore
include an edge
that defines an internal hole in a metal sheet of the alloys described herein,
or an external
edge on such metal sheet. In addition, the edge herein may be formed in a
progressive die
stamping operation which is reference to metal working operation that
typically includes
punching, shearing, coining and bending. The edge herein may be present in a
vehicle, or
more specifically, part of a vehicular frame, vehicular chassis, or vehicle
panel.
Reference to edge expansion herein is understood as increasing the length of
such edge with a
corresponding change in the thickness of the edge. That is confirmed by the
above data in
Table 48, which shows that with respect to an edge that is present in a hole,
when such edge
in the hole is expanded at a speed of greater than or equal to 5 mm/min, one
observes an
increase in the hole expansion ratio (i.e. the edge in the hole is capable of
expansion to higher
percentages over the original diameter) and the edge getting thinner as shown
for example in
the cross sections of the expanded edges in Fig. 91.
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Case Example 20 HER Performance as a Function of Punch Speed and Hole
Expansion
Speed
Sheet from Alloy 9 was produced according to the atomic ratios provided in
Table 2. Slabs
produced by continuous casting were hot rolled into hot band which was
subsequently
processed into sheet with thickness of approximately 1.4 mm by cold rolling
and annealing
cycles. The microstructure of the produced sheet using both SEM and etched
optical
microscopy is demonstrated in FIG. 85 showing typical Recrystallized Modal
Structure.
In Fig. 85A and Fig. 85B, the SEM micrographs shows the micron scale nature of
the
austenitic grains which contain some annealing twins and stacking faults. In
Fig. 85C and
Fig. 85D, etched samples were examined using optical microscopy. It can be
seen that the
grain boundaries are preferentially etched and the microstructure showing the
grain
boundaries. The grain size was measured with the line intercept method and is
found to
range from 6 um to 22 um with a mean value of 15 um.
The sheet with Recrystallized Modal Structure was used for HER testing.
Specimens for
testing with a size of 89 x 89 mm were cut via wire EDM from the sheet. Holes
of 10 mm
diameter were punched at two different speeds of 5 mm/s using an Instron
mechanical test
frame and at 228 mm/s using a commercial punch press with a flat punch
geometry and with
punch to die clearances of approximately 12.5% and 16%, respectively.
Hole expansion ratio (HER) testing was performed on an Interlaken Technologies
SP-225
hydraulic press and consisted of raising the conical punch that uniformly
expanded the hole
radially outward. Two hole expansion speeds of 3 mm/min and 50 mm/min,
synonymous
with the conical ram travel speed, were used. A digital image camera system
was focused on
the conical punch and the edge of the hole was monitored for evidence of crack
formation
and propagation.
The initial diameter of the hole was measured twice with calipers,
measurements were taken
at 90 increments and averaged to get the initial hole diameter. The conical
punch was raised
continuously until a crack was observed propagating through the specimen
thickness. At that
point the test was stopped and the hole expansion ratio was calculated as a
percentage of the
initial hole diameter measured before the start of the test. After expansion
four diameter
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measurements were taken using calipers every 45 and averaged to account for
any
asymmetry of the hole due to cracking.
Hole expansion ratio values for tests are listed in Table 49. HER values vary
from 2.4 to
18.5% in the samples with holes punched at 5 mm/s. In the case of 228 mm/s
hole punching
speed, HER values are significantly higher in a range from 33.8 to 75.0%. The
effect of
expansion speed is illustrated in FIG. 86. Increase in expansion speed results
in higher HER
values independent of utilized punching speeds (i.e. 5 mm/s and 228 mm/s).
Table 49 Hole Expansion Ratio in Alloy 9 Sheet at Different Punching and
Expansion
Speeds
Punch Hole Punch
Hole Expansion
Clearance Speed HER (%)
(%) (mm/s) Speed (mm/min)
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
The magnetic phases volume percent (Fe%) was measured in the HER tested
samples with
different hole punching speed and hole expansion speed using a Fischer
Feritscope FMP30.
The results are listed in Table 50. FIG. 87 illustrates the effect of on the
magnetic phases
volume percent in the tested samples as a function of the distance from the
hole edge. As can
be seen with higher punch speed and/or higher expansion speed, after testing
is completed,
the magnetic phase volume % increases near the hole edge and also away from
the hole edge
into the material. As the increase in magnetic phase volume (Fe%) is
consistent with
increases in the amount of Structure #5 in Table 1 which is formed during
deformation, due
to the formation of magnetic nanoscale alpha-iron from the starting non-
magnetic austenite
present in Structure #4.
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Table 50 Magnetic Phases Volume (Fe%) in Alloy 9 at Different Hole Punching
Speeds
and Hole Expansion Speeds as a Function of Distance from Hole Edge After
Expansion
Hole Creation and Expansion Parameters
Hole Punching Speed
228 228 228 228 5 5 5 5
(mm/s)
Punch Clearance (%) 16 16 16 16 12.5
12.5 12.5 12.5
Hole Expansion Speed
3 3 50 50 3 3 50 50
(mm/min)
HER (%) 33.8 41.3 63.1 75.0 2.4 7.9 12.7
18.5
Distance from hole (mm) Magnetic Phases Volume % (Fe%)
1 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
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 Case Example illustrates that the relative resistance to cracking of an
edge as confirmed
5 by HER testing can be increased by, in the exemplary case of forming an
edge within a hole,
by either increasing hole punching speeds, hole expansion speeds or both. The
sheet from
Alloy 9, tested in this Case Example, demonstrated an increase in hole
expansion ratio with
increasing hole punching speed (i.e. 5 to 228 mm/s) and / or the hole
expansion speed (i.e. 3
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to 50 mm/min). Accordingly, preferably herein for the subject alloys, one
forms an edge in
the alloy and expands the edge at a speed of greater than or equal to 5
mm/min, The
magnetic phases volume percent (Fe%) in tested samples increases with
increasing hole
punching speed and / or the hole expansion speed over the ranges studied. With
this
relatively greater amount of deformation available in and adjacent to the hole
edge during the
now disclosed increased hole punching speed or hole expansion speed, the
higher local
formability and resistance to cracking of the edge is achieved in the material
as measured by
the HER.
Case Example #21 HER Performance as a Function of Hole Preparation Method
Slabs with thickness of 50 mm were laboratory cast from three selected alloys
according to
the atomic ratios provided in Table 2 and laboratory processed by hot rolling,
cold rolling and
annealing at 850 C for 10 min as described herein. Sheet from each alloy
possessing the
Recrystallized Modal Structure with final thickness of 1.2 mm were used to
demonstrate an
effect of hole expansion speed on HER performance.
Specimens for testing with a size of 89 x 89 mm were cut via wire EDM from the
sheet. A
10 mm diameter hole was prepared by various methods including punching, EDM
cutting,
milling, and laser cutting. Hole punching was done at a low quasistatic
punching speed of
0.25 mm/s at 16% punch to die clearance using a Komatsu 0B580-3 press. EDM cut
holes
were first rough cut then the final cut was made at parameters to yield a
visually smooth
surface. During hole milling, holes were pilot drilled, reamed to size, and
then deburred.
Laser cut samples were cut on a 4kW fiber optic Mazak Optiplex 4020 Fiber II
machine.
Hole expansion ratio (HER) testing was performed on an Interlaken Technologies
SP-225
hydraulic press and consisted of raising the conical punch that uniformly
expanded the hole
radially outward. In FIG. 88, the results of HER testing is provided for each
alloy as a
function of the hole preparation method. As shown, in the case of punched
holes, HER
values are the lowest for all three alloys in a range from 6 to 12%. Samples
with EDM cut,
milled and laser cut holes exhibit high HER values from 65 to 140%+. Note that
the ¨140%
expansion represented the maximum extension limit of the press crosshead
during testing so,
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in the samples with EDM cut holes from Alloy 12, and with milled holes from
Alloy 9 and
Alloy 12, the expansion limit was not reached during HER testing (i.e. actual
value > 140%).
In Figure 89, SEM images of the sample cross section near the hole edge prior
to expansion
are provided at low magnification for samples from Alloy 1 with holes prepared
by different
methods. In the punched sample (FIG. 89A), the typical rollover zone at the
top and burr
zone at the bottom can be seen. Additionally, a hemispherical shear affected
zone is visible
at the edge of the hole with penetration of ¨0.5 mm at the deepest point. A
similarly shear
affected zone was observed in punched samples from the other two alloys as
well but not in
any of the samples with holes produced by the non-punching methods. Note that
every
method utilized for hole preparation introduced some kind of defects at the
hole edge. In the
EDM cut hole (FIG. 89B), the edge is perpendicular on a cross section image
but small
micron scaled cutting defects can be seen at the surface; in the milled
samples (FIG. 89C) the
edge of the holes is trapezoidal in shape; and in the laser cut holes (FIG.
89D), the edge
wandered in a sideways fashion as the laser penetrated the sample. In FIG. 90,
SEM images
of the cross sections near the hole edge (i.e. at the edge and up to 0.7 mm
away from the
edge) prior to expansion are provided at higher magnification for samples from
Alloy 1 with
holes prepared by different methods including punching at a hole punching
speed of 0.25
mm/s, EDM cut hole, milled hole, and laser cut hole. The microstructure near
the hole edges
are illustrated in FIGS. 90A, 90B, 90C and 90D respectively. As can be seen,
the edge of the
hole punched at 0.25 mm/s (FIG. 90a) is relatively highly deformed thereby
leading to the
observed low HER values. This structure near the edge of the punched sample is
representative of Structure #5 Refined High Strength Nanomodal Structure in
Table 1
whereby the structures near the hole edge of the EDM cut, milled, and laser
cut holes, is
representative of Structure #4 Recrystallized Modal Structure in Table 1.
However, in
examples where holes were prepared by non-punching methods (FIGS. 90B, 90C,
90D), the
resulting alloys experienced excellent local formability with high HER values
from 65 to
140%+ consistent with the ductile nature of Structure #4 near the hole edges.
In FIG. 91A
(punched hole), 91B (EDM cut hole), 92C (milled hole), and 91D (laser cut
hole), SEM
images of the cross section near the hole edge after HER testing are provided
at low
magnification for samples from Alloy 1. Note that the thickness of samples
near the hole is
smaller in the expanded samples with higher HER values since the expansion of
the holes
results in sample thinning near the hole edge.
CA 03053383 2019-08-12
WO 2018/160387
PCT/US2018/018751
124
In FIG. 92, images of sample cross sections near the hole edge after HER
testing (i.e. after
expansion until failure by cracking) are provided at higher magnification for
samples from
Alloy 1 with holes prepared by different methods showing similar deformed
structure in all
cases. Since hole expansion and deformation of the edge is complete, the
microstructure near
all of the hole edges are similiar and representative of Structure #5 Refined
High Strength
Nanomodal Structure in Table 1.This Case Example demonstrates the effect of
edge
preparation on the resulting local formability in alloys herein. Punching at a
low speeds of
0.25 mm/s is causing structural changes near the hole edge consistent with
previous case
examples resulting in limited local formability of the edges and low HER
values. However,
in examples where holes were prepared by non-punching methods, the resulting
alloys
experienced excellent local formability with high HER values from 65 to 140%+
consistent
with the ductile nature of the microstructure in the samples and at the hole
edges.
