Note: Descriptions are shown in the official language in which they were submitted.
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COLD-WORKED STEELS WITH
PACKET-LATH MARTENSITE/AUSTENITE
MICROSTRUCTURE
BACKGROUND OF THE INVENTION
1. Field of the Invention
[0002] This invention resides in the technology of low and medium carbon steel
alloys, particularly those of high-strength and toughness, and the cold
formability of such
alloys.
2. Description of the Prior Art
[0003] An important step in the processing of high-performance steels is cold
working, which typically consists of a series of compressions and/or
expansions
achieved by processes such as drawing, extruding, cold heading, or rolling.
Cold
working causes plastic deformation of the steel which produces strain
hardening while
forming the steel into the shape in which it will ultimately be used. Cold
working, which
in the case of steel wire is performed by wire drawing, is typically performed
in a
succession of stages with intermediate heat treatments, which in the case of
steel wire are
termed "patenting."
[0004] High-strength steel wire is an example of a high-performance steel and
is
useful in a variety of engineering applications including tire cord, wire
rope, and strand
for pre-stressed concrete reinforcements. The steel most commonly used in high-
strength
steel wire is medium-or high-carbon steel. In the typical procedure for
forming the wire,
hot-rolled rods with pearlitic microstructures are cold drawn in several
stages, with
intermediate patenting treatments to soften the pearlite for continued cold
drawing. For
example, hot rolled rods of about 5. 5 mm diameter might be coarse drawn in
several
stages to a diameter of about 3 mm. Patenting might then be performed at 850-
900 C,
causing austenitization of the steel, followed by transformation of the steel
at 500-550 C
to fine pearlitic lamellae. The steel would then be pickled, in hydrochloric
acid, for
example, to remove the scale formed during patenting. The pickling would be
followed
by several further drawing stages to reduce the diameter down to about 1 mm,
then
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further patenting and pickling. The final drawing would then be done in
several stages to
the final desired diameter, which may for example be about 0.4 mm, to achieve
the
desired properties, notably strength. This may be followed by further
processing such as
stranding, depending on the ultimate use.
[0005] The purpose of the initial patenting treatment is to produce a wire rod
with a
fine lamellar pearlite structure, which requires a low transformation
temperature. To
achieve the desired temperature control, the process is typically performed in
a molten
lead bath. In the succeeding drawing stages, the wire is drawn to true strains
(defined
below) of 6-7 to obtain high strength levels of approximately 3,000 MPa. For
conventional pearlitic wires, these high strains and strengths are attainable
only by
applying a series of patenting treatments. Without these patenting treatments,
the cold
drawing will cause shear cracking of the pearlitic lamellae. Because of the
need for a
molten lead bath the entire process is costly and tends to raise environmental
concerns.
[0006] Cold working is also used in the production of expandable steel tubing,
i. e. ,
tubing that is expanded on-site and in some cases below ground.
[0007] A recent development in steel alloys is the formation of
microstructures
containing both martensite and austenite phases in an alternating
configuration in which
the martensite is present as laths that are separated by thin films of
austenite. The
microstructures are fused grains in which individual grains contain several
laths of
martensite separated by thin austenite films with, in some cases, an austenite
shell
surrounding each grain. These structures are termed" dislocated lath
martensite"structures
or"packet-lath" martensite/austenite"structures. Patents disclosing these
microstructures
are as follows:
4,170, 497 (Gareth Thomas and Bangaru V. N. Rao), issued October 9,1979 on
an application filed August 24,1977
4,170, 499 (Gareth Thomas and Bangaru V. N. Rao), issued October 9,1979 on
an application filed September 14,1978 as a continuation-in-part of the
above application filed on August 24,1977
4,671, 827 (Gareth Thomas, Nack J. Kim, and Ramamoorthy Raines), issued
June 9,1987 on an application filed on October 11,1985
6,273, 968 B 1 (Gareth Thomas), issued August 14,2001 on an application filed
on March 28,2000
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While these microstructures offer certain performance benefits, notably a high
resistance
to corrosion, it has not heretofore been known that processing steps typically
used for
steel alloys could be simplified or eliminated when these microstructures are
present.
[00081 Of further potential relevance to this invention are two United States
patents
that disclose the cold working of steel rods and wires without patenting.
These patents
are:
4,613, 385 (Gareth Thomas and Alvin H. Nakagawa), issued September 23, 1986
on an application filed December 9,1982
4,619, 714 (Gareth Thomas, Jae-Hwan Ahn, and Nack-Joon Kim), issued
October 28,1986 on an application filed November 29,1984 as a
continuation-in-part of the above application filed on August 6,1984
The microstructures of the steels in these patents are considerably different
from those of
the first four patents listed above.
SUMMARY OF THE INVENTION
[00091 It has now been discovered that the packet-lath martensite/ austenite
microstructure is unique in its crystallographic characteristics and how these
characteristics cause it to respond to cold working. Because of the high
dislocation
density of this microstructure and the ease with which strains in the
structure can move
between the martensite and austenite phases, cold working provides the
microstructure
with unique mechanical properties that include a high tensile strength. As a
result, these
alloys can be cold worked without intermediate heat treatments, while still
achieving
tensile strengths comparable to the tensile strengths of conventional steel
alloys that have
been processed by cold working with intermediate heat treatments. In the case
of steel
wire having the packet-lath martensite/austenite microstructure, this
invention lies in the
discovery that cold drawing can be performed without intermediate patenting
treatments.
In accordance with the present invention, therefore, carbon steel alloys
having the
packet-lath martensite/austenite microstructure, i. e. , those whose
microstructure
includes laths of martensite alternating with thin films of retained austenite
are cold
formed, preferably without intermediate heat treatments, to a reduction
sufficient to
achieve a tensile strength of about 150 ksi or higher ("ksi"denotes kilo-
pounds-force per
square inch), equivalent to approximately 1,085 MPa or higher ("MPa"denotes
megapascals, i. e. , newtons per square millimeter). Cold working to tensile
strengths of
2,000 MPa (290 ksi) of higher is of particular interest, and indeed, tensile
strengths of
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3,000 MPa (435 ksi) and as high as 4,000 MPa (580 ksi) can be achieved by the
practice of this
invention. These values are approximate; the conversion factor to the nearest
thousandth is 6.895
MPa equal 1 ksi.
[0010] The benefits of this invention extend to simple packet-lath
martensite/austenite
microstructures containing no ferrite or insignificant amounts of ferrite, and
also to
microstructures that include packet-lath grains fused with ferrite grains, and
to variants on these
structures, including those whose packet-lath grains are encased by austenite
shells, those that are
free of interphase carbide precipitates, and those in which the austenite
films are of a uniform
orientation. The discovery of the ability of packet-lath martensite/austenite
microstructures to
respond to cold working in this manner is surprising relative to the
disclosures in patents nos.
4,613, 385 and 4,619, 714 referenced above, since the ferrite in the
microstructures of those
patents has a lower yield strength than the martensite. As a result, the
ferrite will preferentially
absorb the strain introduced by the cold working, while the martensite will
not respond to the cold
working until the ferrite phase is work hardened to a level above the yield
strength of the
martensite. In the microstructures addressed by the present invention, the
relatively low level of
ferrite, or its absence when no ferrite is present, will cause the martensite
to absorb the strain at an
earlier stage of the cold working process.
Martensite and ferrite are distinctly different from each other in crystal
structure and hardening
behavior.
[0010a] In accordance with an illustrative embodiment, there is provided a
process for
manufacturing a high-strength, high-ductility alloy carbon steel, said process
comprising: (a)
forming a carbon steel alloy having a microstructure comprising laths of
martensite alternating
with films of retained austenite, and (b) cold working said carbon steel alloy
in a series of passes
without heat treatment between passes to a reduction sufficient to achieve a
tensile strength of at
least about 150 ksi, wherein said alloy carbon steel contains one of: from
0.04% to 0.12% carbon,
from zero to 11% chromium, from zero to 2.0% manganese, and from zero to 2.0%
silicon, all by
weight, the remainder being iron together with any unavoidable impurities, and
from 0.02% to
0.14% carbon, from zero to 3.0% silicon, from zero to 1.5% manganese, and from
zero to 1.5%
aluminum, all by weight, the remainder being iron together with any
unavoidable impurities.
100111 These and other features, objects, advantages, and embodiments of the
invention will
be better understood from the descriptions that follow.
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BRIEF DESCRIPTION OF THE FIGURES
[0012] FIG. 1 is a plot of tensile strength vs. true total strain for two
steel alloys of dual-
phase packet-lath martensite/austenite microstructure, upon cold working in
accordance with
this invention in the absence of intermediate heat treatments.
[0013] FIG. 2 is a plot of tensile strength vs. true total strain for three
steel alloys of triple-
phase packet-lath martensite/austenite/ferrite microstructure and one steel
alloy of dual-phase
packet-lath martensite/austenite microstructure, upon cold working in
accordance with this
invention in the absence of intermediate heat treatments.
DETAILED DESCRIPTION OF THE INVENTION
AND PREFERRED EMBODIMENTS
[0014] Cold working in the practice of this invention can be performed by the
use of
techniques and equipment that have been used for cold working in the prior art
on other steel
alloys and microstructures. For alloys in the form of blooms, billets, bars,
slabs or sheets,
cold working may consist of rolling the steel between rollers or other means
of compression
to reduce the thickness of and elongate the steel. When cold working is
performed by rolling,
multiple reductions are achieved by multiple passes through a rolling mill.
For rod-shaped or
wire-shaped workpieces, cold working may consist of cold-drawing or extrusion
through a
die. For multiple reductions, the workpiece is extruded through a series of
successively
smaller dies. Tubing is achieved by drawing the steel through a ring-shaped
die with a
mandrel inside the die. For multiple passes, the tubing that has already been
drawn is further
drawn through a smaller ring-shaped die with a mandrel placed inside the
tubing.
[0015] Cold working is performed at a temperature below the lowest temperature
at which
recrystallization occurs. Suitable temperatures are therefore those that do
not induce any
phase change in the steel. For carbon steels, recrystallization typically
occurs at
approximately 1,000 C (1,832 F), and accordingly, cold working in accordance
with this
invention is performed well below this temperature. Preferably, cold working
is performed at
temperatures of about 500 C (932 F) or less, more preferably about 100 C (212
F) or less,
and most preferably at a temperature that is within about 25 C of ambient
temperature.
[0016] Cold working can be performed in a single pass or in a succession of
passes. In
either case, intermediate heat treatments (which, in the case of steel wire,
are termed
"patenting") may be performed for further improvement in properties, but the
properties
resulting from the cold working alone are sufficiently high that the
intermediate heat
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treatments are not required and are preferably not performed. The degree of
reduction per
pass is not critical to the invention and can vary widely, although the
reductions should
be great enough to avoid hardening the steel so much that the steel becomes
susceptible
to breakage after a small total reduction. In most cases, preferred reductions
are at least
about 20% per pass, more preferably at least about 25% per pass, and most
preferably
from about 25% to about 50% per pass. The reduction per pass is at least
partially
governed by such factors as the die angle and the drawing efficiency
coefficient. The
larger the die angle, the larger the minimum reduction that is required to
avoid central
burst cracking. The lower the drawing efficiency coefficient, however, the
lower the
maximum reduction for a steel with a given strain hardening exponent. A
compromise is
typically sought between these two competing considerations. In terms of the
tensile
strength of the final product, the cold working will preferably be performed
to a tensile
strength within the range of from about 150 ksi to about 500 ksi.
[00171 The process of this invention is applicable to carbon steel alloys
having
packet-lath martensite/austenite microstructures such as those described in
the patents
cited above, as well as those described in co-pending United States Patent
Applications
Nos. 10/017, 847, filed December 15,2001 (entitled "Triple-Phase Nano-
Composite
Steels, "inventors Kusinski, G. J. , Pollack, D. , and Thomas, G. ), and
10/017, 879, filed
December 14,2001 (entitled "Nano-Composite Martensitic Steels, "inventors
Kusinski,
G. J. , Pollack, D. , and Thomas, G. ). To permit formation of the packet-lath
martensite/austenite microstructure, the alloy composition will typically have
a
martensite start temperature MS of about 300 C or higher, and preferably 350
C or
higher. While alloying elements in general affect the Ms, the alloying element
that has
the strongest influence on the MS is carbon, and achieving an alloy with MS
above 300 C
can be achieved by limiting the carbon content of the alloy to a maximum of
0.35% by
weight. In preferred embodiments of the invention, the carbon content is
within the range
of from about 0.03% to about 0.35%, and in more preferred embodiments, the
range is
from about 0.05% to about 0.33%, all by weight. Further alloying elements,
such as
molybdenum, titanium, niobium, and aluminum, can also be present in amounts
sufficient to serve as nucleation sites for fine grain formation yet low
enough in
concentration to avoid affecting the properties of the finished alloy by their
presence.
The concentration should also be low enough to avoid the formation of
inclusions and
other large precipitates, which may render the steel susceptible to early
fracture. In
certain embodiments of the invention, it will be
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advantageous to include one or more austenite stabilizing elements, examples
of which are
nitrogen, manganese, nickel, copper, and zinc. Particularly preferred among
these are
manganese and nickel. When nickel is present, the nickel concentration is
preferably within
the range of about 0.25% to about 5%, and when manganese is present, the
manganese
concentration is preferably within the range of from about 0.25% to about 6%.
Chromium is
also included in many embodiments of the invention, and when it is present,
the chromium
concentration is preferably from about 0.5% to about 12%. All concentrations
herein are by
weight.
[0018] Certain embodiments of the invention involve alloys that include a
ferrite phase in
addition to the packet-lath martensite/austenite grains (triple-phase alloys)
while others
contain only the packet-lath martensite/austenite grains and do not include a
ferrite phase
(dual-phase alloys). In general, the presence or absence of the ferrite phase
is determined by
the type of heat treatment in the initial austenitization stage. By
appropriate selection of the
temperature, the steel can be transformed into a single austenite phase or
into a two-phase
structure containing both austenite and ferrite. In addition, the alloy
composition can be
selected or adjusted to either cause ferrite formation during the initial
cooling of the alloy
from the austenite phase or to avoid ferrite formation during the cooling,
i.e., to avoid the
formation of ferrite grains prior to the further cooling of the austenite to
form the packet-lath
microstructure.
[00191 As noted above, in certain cases it will be beneficial to use alloys
with packet-lath
martensite/austenite microstructures in which the austenite films in a single
packet-lath grain
are all of approximately the same orientation, although the crystallographic
orientation may
vary, or those in which the austenite films in a single packet-lath grain are
all of the same
crystal plane orientation. The latter can be achieved by limiting the grain
size to ten microns
or less. Preferably, the grain size in these cases is within the range of
about 1 micron to about
10 microns, and most preferably from about 5 microns to about 9 microns.
[00201 The preparation of -phase packet-lath martensite/austenite
microstructures that do
not contain ferrite (i.e., "dual-phase" microstructures) begins with the
selection of the alloy
components and the combining of these components in the appropriate portions
as indicated
above. The combined components are then homogenized ("soaked") for a
sufficient period
of time and at a sufficient temperature to achieve a uniform austenitic
structure with all
elements and components in solid solution. The temperature will be above the
austenite
recrystallization temperature but preferably at a level that will cause very
fine grains to form.
The austenite recrystallization temperature typically varies with the alloy
composition, but in
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general will be readily apparent to those skilled in the art. In most cases,
best results will be
achieved by soaking at a temperature within the range of 800 C to 1150 C.
Rolling, forging
or both are optionally performed on the alloy at this temperature.
[0021] Once homogenization is completed, the alloy is subjected to a
combination of
cooling and grain refinement to the desired grain size, which as noted above
may vary. Grain
refinement may be performed in stages, but the final grain refinement is
generally achieved at
an intermediate temperature that is above, yet close to, the austenite
recrystallization
temperature. The alloy may first be rolled at the homogenization temperature
to achieve
dynamic recrystallization, then cooled to an intermediate temperature and
rolled again for
further dynamic recrystallization. The intermediate temperature is between the
austenite
recrystallization temperature and a temperature that is about 50 degrees
Celsius above the
austenite recrystallization temperature. For alloy compositions whose
austenite
recrystallization temperature is about 900 C, and the intermediate temperature
to which the
alloy is cooled is preferably between about 900 to about 950 C, and most
preferably
between about 900 to about 925 C. For alloy compositions whose austenite
recrystallization
temperature is about 820 C, the preferred intermediate temperature is about
850 C. Dynamic
recrystallization can also be achieved by forging or by other means known to
those skilled in
the art. Dynamic recrystallization produces a grain size reduction of 10% or
greater, and in
many cases a grain size reduction of from about 30% to about 90%.
[0022] Once the desired grain size is achieved, the alloy is quenched by
cooling from a
temperature above the austenite recrystallization temperature down to the
martensite start
temperature MS, then through the martensite transition range to convert the
austenite crystals
to the packet-lath martensite/austenite microstructure. When ferrite crystals
are present
among the austenite crystals, the conversion occurs only in the austenite
crystals. The
optimal cooling rate varies with the chemical composition, and hence the
hardenability, of the
alloy. The resulting packets are of approximately the same small size as the
austenite grains
produced during the rolling stages, but the only austenite remaining in these
grains is in the
thin films and in some cases in the shell surrounding each packet-lath grain.
When the thin
austenite films are to be of a single variant in crystal orientation, this is
achieved by
controlling the process to achieve a grain size of less than 50 microns.
[0023] As an alternative to dynamic recrystallization, grain refinement to the
desired grain
size can be accomplished by heat treatment alone. To use this method, the
alloy is quenched
as described in the preceding paragraph, then reheated to a temperature that
is approximately
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equal to the austenite recrystallization temperature or slightly below, then
quenched once
again to achieve, or to return to, the packet-lath martensite/austenite
microstructure. The
reheating temperature is preferably within about 50 degrees Celsius of the
austenite
recrystallization temperature, for example about 870 C.
[00241 Processing steps such as heating the alloy composition to the austenite
phase,
cooling the alloy with controlled rolling or forging to achieve the desired
reduction and grain
size, and quenching the austenite grains through the martensite transition
region to achieve
the packet-lath structure are performed by methods known in the art. These
methods include
castings, heat treatment, and hot working of the alloy such as by forging or
rolling, followed
by finishing at the controlled temperature for optimum grain refinement.
Controlled rolling
serves various functions, including aiding in the diffusion of the alloying
elements to form a
homogeneous austenite crystalline phase and in the storage of strain energy in
the grains. In
the quenching stages of the process, controlled rolling guides the newly
forming martensite
phase into a packet-lath arrangement of martensite laths separated by thin
films of retained
austenite. The degree of rolling reduction can vary and will be readily
apparent to those of
skill in the art. Quenching is preferably done fast enough to avoid formation
of detrimental
microstructures including pearlite, bainite, and particles or precipitates,
particularly
interphase precipitation and particle formation, including the formation of
undesirable
carbides and carbonitrides. In the packet-lath martensite-austenite grains,
the retained
austenite films will constitute from about 0.5% to about 15% by volume of the
microstructure, preferably from about 3% to about 10%, and most preferably a
maximum of
about 5%.
[00251 Triple-phase alloys have a microstructure consisting of two types of
grains, ferrite
grains and packet-lath martensite/austenite grains, fused together as a
continuous mass. As in
the dual-phase alloys, the individual grain size is not critical and can vary
widely. For best
results, the grain sizes will generally have diameters (or other appropriately
characteristic
linear dimension) that fall within the range of about 2 microns to about 100
microns, or
preferably within the range of about 5 microns to about 30 microns. The amount
of ferrite
phase relative to the martensite-austenite phase may vary. In most cases,
however, best
results will be obtained when the martensite/austenite grains constitute from
about 5% to
about 95% of the triple-phase structure, preferably from about 15% to about
60%, and most
preferably from about 20% to about 40%, all by weight.
[00261 Triple-phase alloys can be prepared by first combining the appropriate
components
needed to form an alloy of the desired composition, then soaking to achieve a
uniform
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austenitic structure with all elements and components in solid solution, as in
the preparation
of the dual-phase alloys described above. A preferred soaking temperature
range is from
about 900 C to about 1,170 C. Once the austenite phase is formed, the alloy
composition is
cooled to a temperature in the intercritical region, which is defined as the
region in which
austenite and ferrite phases coexist at equilibrium. The cooling thus causes a
portion of the
austenite to transform into ferrite grains, leaving the remainder as
austenite. The relative
amounts of each of the two phases at equilibrium varies with the temperature
to which the
composition is cooled in this stage, and also with the levels of the alloying
elements. The
distribution of the carbon between the two phases (again at equilibrium) also
varies with the
temperature. The relative amounts of the two phases are not critical to the
invention and can
vary. The temperature to which the composition is cooled in order to achieve
the dual-phase
ferrite-austenite structure is preferably within the range of from about 800 C
to about
1,000 C.
[0027] Once the ferrite and austenite crystals are formed (i.e., once
equilibrium at the
selected temperature in the intercritical phase is achieved), the alloy is
rapidly quenched by
cooling through the martensite transition range to convert the austenite
crystals to the packet-
lath martensite/austenite microstructure. The cooling rate used during this
transition is great
enough to substantially avoid any changes to the ferrite phase and to avoid
undesirable
austenite decomposition. Depending on the alloy composition and its
hardenability, water
cooling may be required to achieve the desired cooling rate, although for
certain alloys air
cooling will suffice. In some alloys, notably triple-phase containing 6% Cr,
the desired
cooling rate is slow enough that air cooling can be used. The considerations
noted above in
connection with dual-phase alloys apply here as well.
[0028] Preferred dual-phase alloy compositions are those that contain from
about 0.04% to
about 0.12% carbon, from zero to about 11.0% chromium, from zero to about 2.0%
manganese, and from zero to about 2.0% silicon, all by weight, the remainder
being iron.
Preferred triple-phase alloy compositions are those that contain from about
0.02% to about
0.14% carbon, from zero to about 3.0% silicon, from zero to about 1.5%
manganese, and
from zero to about 1.5% aluminum, all by weight, the remainder being iron.
[0029] The formation of precipitates or other small particles within the
microstructure upon
cooling is collectively referred to as "autotempering." In certain
applications of this
invention, whether dual-phase or triple-phase alloys, autotempering will
purposely be
avoided by using a relatively fast cooling rate. The minimum cooling rates
that will avoid
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autotempering are evident from the transformation-temperature-time diagram for
the alloy.
In the typical diagram, the vertical axis represents temperature and the
horizontal axis
represents time, while curves on the diagram indicate the regions where each
phase exists
either by itself or in combination with another phase(s). A typical such
diagram is shown in
Thomas, U.S. Patent No. 6,273,968 B1, referenced above. In such diagrams, the
minimum
cooling rate is a line of descending temperature over time which abuts the
left side of a C-
shaped curve. The region to the right of the curve represents the presence of
carbides, and
cooling rates that avoid carbide formation are therefore those represented by
lines that remain
to the left of the curve. The line that is tangential to the curve has the
smallest slope and is
therefore the slowest rate that can be used while still avoiding carbide
formation.
[0030] The terms "interphase precipitation" and "interphase precipitates" are
used herein to
denote the formation of small alloy particles at locations between the
martensite and austenite
phases, i.e., between the laths and the thin films separating the laths.
"Interphase
precipitates" does not refer to the austenite films themselves. Interphase
precipitates are to be
distinguished from "intraphase precipitates," which are precipitates located
within the
martensite laths rather than along the interfaces between the martensite laths
and the austenite
films. Intraphase precipitates that are about 500A or less in diameter are not
detrimental to
toughness and may in fact enhance toughness. Thus, autotempering is not
necessarily
detrimental provided that the autotempering is limited to intraphase
precipitation and does not
result in interphase precipitation. The term "substantially no carbides" is
used herein to
indicate that if any carbides are present, their distribution and amount are
such that they have
a negligible effect on the performance characteristics, and particularly the
corrosion
characteristics, of the finished alloy.
[0031] Depending on the alloy composition, a cooling rate that is sufficiently
high to
prevent carbide formation or autotempering in general may be one that can be
achieved with
air cooling or one that requires water cooling. In alloy compositions in which
autotempering
can be avoided with air cooling, air cooling can still be done when the levels
of certain
alloying elements are reduced provided that the levels of other alloying
elements are raised.
For example, a reduction in the amount of carbon, chromium, or silicon can be
compensated
for by raising the level of manganese.
[0032] The processes and conditions set forth in the U.S. patents referenced
above,
particularly heat treatments, grain refinements, on-line forgings and the use
of rolling mills
for rounds, flats, and other shapes, may be used in the practice of the
present invention for the
heating of the alloy composition to the austenite phase, the cooling of the
alloy from the
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austenite phase to the intercritical phase in the case of triple-phase alloys,
and then the
cooling through the martensite transition region. Rolling is performed in a
controlled manner
at one or more stages during the austenitization and first-stage cooling
procedures, for
example, to aid in the diffusion of the alloying elements to form a
homogeneous austenite
crystalline phase and then to deform the crystal grains and store strain
energy in the grains,
while in the second-stage cooling, rolling can serve to guide the newly
forming martensite
phase into the packet-lath arrangement of martensite laths separated by thin
films of retained
austenite. The degree of rolling reductions can vary, and will be readily
apparent to those
skilled in the art. In the packet-lath martensite-austenite crystals, the
retained austenite films
will constitute from about 0.5% to about 15% by volume of the microstructure,
preferably
from about 3% to about 10%, and most preferably a maximum of about 5%. The
proportion
of austenite relative to the entire triple-phase microstructure will be a
maximum of about 5%.
The actual width of a single retained austenite film is preferably within the
range of about
50A to about 250A, and preferably about 100A. The proportion of austenite
relative to the
entire triple-phase microstructure will in general be a maximum of about 5%.
The rolling
discussed in this paragraph is to be distinguished from the cold working that
is done in
accordance with this invention after the packet-lath martensite/austenite
microstructures,
whether dual-phase or part of a triple-phase structure, have been formed.
[0033] The following examples are offered only by way of illustration.
EXAMPLE 1
[0034] This example illustrates the deformation of a carbon steel rod with a
packet-lath
martensite/austenite microstructure, by a cold drawing process according to
the present
invention to an area reduction of 99%.
[0035] The experiment reported in this example was performed on a steel rod
measuring
6 mm in diameter and having an alloy composition of 0.1% carbon, 2.0% silicon,
0.5%
chromium, 0.5% manganese, all by weight, and the balance iron, with a
microstructure
consisting of grains measuring approximately 50 microns in diameter, each
grain consisting
of laths of martensite measuring approximately 100 nm in thickness alternating
with thin
films of austenite measuring approximately 10 nm in thickness, with no ferrite
phases and
each grain surrounded by an austenite shell measuring approximately 10 nm in
thickness.
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The rod was prepared by the method described in co-pending United States
patent application
serial no. 10/017,879, filed December 14, 2001, referenced above.
[0036] The uncoated steel rod was surface cleaned and lubricated, then cold
drawn through
lubricated dies in 15 passes at a temperature of 25 C to a diameter of 0.0095
inch (0.024 cm).
At a final wire diameter of 0.0105 inch (0.027 cm), representing a total area
reduction of
99%, the wire had a tensile strength of 390 ksi (2,690 MPa).
EXAMPLE 2
[0037] This example is another illustration of the cold working of carbon
steel rods with
packet-lath martensite/austenite microstructures in accordance with the
present invention. In
this example, two different alloys were used, Fe/8Cr/0.05C and Fe/2Si/0.1C,
with a
microstructure consisting of grains measuring approximately 50 microns in
diameter, each
grain consisting of laths of martensite measuring approximately 150 nm in
thickness
alternating with thin films of austenite measuring approximately 10 nm in
thickness, with no
significant ferrite phases, each grain surrounded by an austenite shell
measuring
approximately 10 nm in thickness.
[0038] The steel rods were 6 mm in diameter, and were surface cleaned and
lubricated,
then cold drawn through lubricated dies in a series of passes at a temperature
of 25 C. The
drawing schedule shown in Table I was used for the Fe/8Cr/0.050 alloy, and a
similar
drawing schedule was used for the Fe/2Si/0.1C alloy. In this table, A
represents the initial
rod diameter and A is the rod diameter after the particular pass.
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TABLE I
Drawing Schedule for Fe/8Cr/0.05C With Substantially
Ferrite-Free Packet-Lath Martensite Microstructure
Single Pass Total
Diameter True Total Strain Area Reduction Area Reduction
Pass No. (mm) (ln(A/A )) (%) (%)
(initial) 6.000 0.0 0.0 0.0
1 4.3 0.7 48.2 48.2
2 3.4 1.1 37.0 67.3
3 2.7 1.6 37.1 79.4
4 2.2 2.0 34.0 86.4
1.8 2.5 36.6 91.4
6 1.4 2.9 38.5 94.7
7 1.0 3.5 45.4 97.1
[0039] Tensile strengths were measured on the starting rod and after each
pass, and the
5 results are plotted against the true total strain in FIG. 1, in which the
squares represent the
Fe/8Cr/0.05C alloy and the diamonds represent the Fe/2Si/0.1C alloy. The
Figure shows that
the tensile strengths of both alloys reach approximately 2,000 MPa by the end
of the entire
drawing sequence at a total area reduction of 97%.
EXAMPLE 3
[0040] This example illustrates cold working in accordance with the present
invention,
using carbon steel rods with packet-lath martensite/austenite microstructures
that contain
ferrite crystals as a third phase (in addition to the laths of martensite and
the thin films of
austenite, i.e., a triple-phase microstructure).
[0041] In this example, the alloy was Fe/2Si/0.1C, with a microstructure
consisting of
ferrite fused with packet-lath grains similar to those described above in
Examples 1 and 2,
containing martensite laths alternating with thin films of austenite and
encased in an austenite
shell. The rods were prepared by the method described in United States patent
application
no. 10/017,847, filed December 14, 2001, referenced above, using a reheat
temperature of
950 C to achieve a ferrite content of 70 volume percent of the microstructure.
The initial rod
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diameter was 0.220 inch (5.59 mm), and the cold working consisted of drawing
the rods
through lubricated conical dies at a temperature of 25 C in 15 passes with
approximately
36% reduction per pass to a final diameter of 0.037 inch (0.94 mm).
[00421 The drawing schedule is shown in Table II, where A represents the
initial rod
diameter and A is the rod diameter after the particular pass.
TABLE II
Drawing Schedule for Fe/2Cr/0.1 C With
Triple-Phase Microstructure
Single Pass Total
Diameter True Total Strain Area Reduction Area Reduction
Pass No. (mm) (ln(A/Ao)) (%) (%)
(initial) 6.050 0.00 0.00 0.00
1 4.580 0.56 42.69 42.69
2 3.650 1.01 36.49 63.60
3 2.910 1.46 36.44 76.86
4 2.320 1.92 36.44 85.29
5 1.870 2.35 35.03 90.45
6 1.660 2.59 21.20 92.47
7 1.320 3.04 36.77 95.24
8 1.090 3.43 31.81 96.75
9 0.910 3.79 30.30 97.74
0.756 4.16 30.98 98.44
11 0.624 4.54 31.87 98.94
12 0.526 4.89 28.94 99.24
13 0.437 5.26 30.98 99.48
14 0.390 5.48 20.35 99.58
0.359 5.65 15.27 99.65
[00431 The tensile strength of the final wire was 2760 MPa (400 ksi).
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EXAMPLE 4
[0044] This example is a further illustration of the cold work of carbon steel
rods whose
microstructure consists of packet-lath martensite/austenite and ferrite
crystals, in accordance
with the present invention.
[0045] In this example, the alloy was Fe/2Si/0.1 C as in Example 3, with a
microstructure
consisting of ferrite fused with packet-lath grains similar to those described
above in
Examples 1 and 2, containing martensite laths alternating with thin films of
austenite and
encased in an austenite shell. A rod of this composition was prepared by the
general method
described in United States patent application no. 10/017,847, filed December
14, 2001,
referenced above. In this case, the rod was initially hot rolled to a diameter
of 0.25 inch
(6.35 mm), then heated to 1,150 C for about 30 minutes to austenitize the
composition, then
quenched in iced brine to transform the austenite to substantially 100%
martensite, then
rapidly reheated to convert the structure to approximately 70% ferrite and 30%
austenite.
The rod was then quenched in iced brine to convert the austenite to the packet-
lath
martensite/austenite structure. The rod was then cold drawn in 7 passes at a
reduction of 35%
per pass to a final diameter of 0.055 inch (1.40 mm), resulting in a tensile
strength of
1,875 MPa (272 ksi). In a parallel experiment, a rod of the same composition
and treated in
the identical manner was cold drawn in 13 passes at a reduction of 35% per
pass to a final
diameter of 0.015 inch (0.37 mm), resulting in a tensile strength of 2,480 MPa
(360 ksi).
EXAMPLE 5
[0046] This example is a still further illustration of the cold working of
carbon steel rods
whose microstructure consists of packet-lath martensite/austenite and ferrite
crystals, in
accordance with the present invention, demonstrating the effect of varying the
relative
amounts of packet-lath martensite/austenite and ferrite.
[0047] The steel alloy was Fe/2Si/0.1C as in Examples 3 and 4, and the rods
were prepared
as described in Example 4, using different reheat temperatures to achieve
ferrite contents of
0%, 56%, 66%, and 75%, corresponding to contents of packet-lath
martensite/austenite
contents of 100%, 44%, 35%, and 25%, respectively, all by volume. Drawing
schedules
similar to that shown in Table II were used on all four microstructures, and
the resulting
tensile strengths are plotted against the true total strain in FIG. 2, in
which the squares
represent the 100% packet-lath alloy, the triangles represent the 44% packet-
lath alloy, the
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circles represent the 34% packet-lath alloy, and the diamonds represent the
25% packet-lath
alloy. The plot shows that all four microstructures achieved a tensile
strength well in excess
of 2,000 MPa, and those in which the packet-lath martensite/austenite portions
exceeded 25%
produced higher tensile strengths than the microstructure in which the packet-
lath portion was
25%.
[0048] The foregoing is offered primarily for purposes of illustration.
Further
modifications and variations of the various parameters of the alloy
composition and the
processing procedures and conditions may be made that still embody the basic
and novel
concepts of this invention. These will readily occur to those skilled in the
art and are
included within the scope of this invention.
17
