Note: Descriptions are shown in the official language in which they were submitted.
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High Nitrogen Steel Powder and Methods of Making Same
CROSS-REFERENCE TO RELATED APPLICATION
100011 This
application claims priority to co-pending U.S. Provisional Patent
Application No. 62/810,680, filed February 26, 2019, the entire contents of
which is hereby
incorporated by reference in its entirety including the drawings.
TECHNICAL FIELD
100011 This disclosure relates to steel powders having dissolved nitrogen in
excess of the
solubility limit at solidification temperature and methods for making such
powders with
tailored phase constituents.
BACKGROUND
100021 Nitrogen (N) is effective in improving the mechanical, wear and
corrosion properties
of steels if it remains in solid solution, specifically in the form of
coherent Cr-N short range
order (SRO). These steels are known as high nitrogen steels (HNS), comprising
of significant
amount of dissolved N (-0.4-1.0 weight percent (wt%)), along with other
alloying elements
such as Cr and Mn and typically containing low nickel or no nickel. Austenitic
stainless steels
with dissolved nitrogen contents up to 0.60 wt% have successfully been
utilized in applications
involving pitting corrosion, crevice corrosion and stress corrosion cracking
in hot chloride
solutions, such as NaCl and MgCl2. However, precipitation of nitride particles
(e.g., CriN, TiN,
VN) leads to Cr and N depletion from the matrix, impairing the corrosion and
wear resistance
of the steel and should be avoided during processing.
100031 The solubility of nitrogen in molten steels at atmospheric pressure is
very low (0.045
wt% at 1600 degrees Celsius ( C)). Normal steel making practice at atmospheric
pressure does
not permit dissolving high amounts of nitrogen into the melt and much of that
dissolved
nitrogen in the liquid is lost during liquid to 8 ferrite solidification due
to even lower N
solubility in 8 ferrite. One way to retain the dissolved nitrogen during
liquid to solid
transformation (or solidification) is by suppressing a liquid (L) ¨>8 reaction
and promoting a
L --4y reaction due to high solubility of N in y phase by alloying the steel
with y stabilizers such
as Ni, Mn and N itself. The development of pressurized metallurgy, namely
rneltina and
solidifying under high N partial pressure, makes it possible to effectively
enhance the dissolved
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N content without inducing nitrogen pores and utilize the beneficial role of
dissolved nitrogen.
But in general these procedures are very expensive and require sophisticated
equipment.
100041 Besides fabricating bulk high nitrogen steel (FINS) by pressurized
metallurgy described
above, powders of }INS can find many beneficial applications such as coatings
and sintered
metal products due to their excellent mechanical and corrosion properties.
Steel powders are
commonly fabricated by atomization of liquid steel. To produce FINS powder by
an
atomization technique, the melting procedure needs to use a high nitrogen
partial pressure
environment along with appropriate alloy composition to avoid an L ¨>E=
reaction and then
atomizing the liquid with high pressure nitrogen gas jets. While this
procedure is feasible, it is
complicated and expensive. Alternatively, mechanical alloying by attrition
milling can
introduce high levels of nitrogen into powders; however, typical time
requirements are in
excess of 100 hours and only a limited amount of material can be processed at
one time.
Further, the process introduces undesirable impurities and hard powders. This
is especially true
ter fabricating coatings using 1-INS powder, where there are not many good
fabrication options
available. These powders cannot be remelted using high temperature coating
processes such as
plasma spray as the dissolved nitrogen will be lost unless high nitrogen
partial pressure is
maintained, during deposition process. Cold spray processes that employ solid
state fusion to
fabricate coatings can be effective in this case, however, it requires that
the powders possess
sufficient deformability (plasticity) that necessitates well controlled
nitrogen dissolution and
phase constituents, preferably y phase free of Cr2N precipitates.
[0005] Providing a means to economically fabricate steel powders having high
level of
dissolved nitrogen without a nitride or oxide compound layer or damaging
brittle nitride
precipitates would benefit many industrial applications where a combination of
high toughness,
wear and corrosion properties are desirable.
SUMMARY
[0006] Provided are methods for the production of steel powders containing
high dissolved
nitrogen, in particular micron size un-sintered powders substantially free of
brittle nitride
compounds.
[0007] Accordingly, a method for producing high nitrogen steel (HNS)
optionally in a powder
form and from a powder precursor is provided, the precursor optionally
comprising, ferrite (a)
phase, or austenite (7) phase or a mixture of a + 7 phase; and the HNS powder
comprising a
mechanically tough alloy having dissolved nitrogen and optionally having a
substantially
homogeneous composition, in weight percent, of from 0.1 to 6.0 wt% nitrogen.
Further, the
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HNS powder may optionally include a single phase nitrogen alloy, optionally a
7 phase alloy.
Thus, an HNS powder is provided comprising a dissolved nitrogen content
significantly in
excess of what could have been achieved through atomizing liquid steel at
atmospheric
pressures.
[0008] Further, an object of the present disclosure is to provide
methodologies to remove
preexisting oxides from the precursor steel powder to accelerate diffusion of
nitrogen into the
powder and prevent precipitation of incoherent nitride precipitates to promote
plastic
deformability. Methodologies as provided herein include exposing the precursor
steel powder
to a reducing gas environment at elevated temperatures and hydrogen from the
reducing gas
environment combines with the preexisting oxygen of the precursor, resulting
in a volatile
byproduct which is removed from the atmosphere, and further quenching the
powders quickly
to ambient temperatures after nitrogen dissolution to prevent precipitation of
nitrides or
formation of oxides. The oxygen removal methods optionally further include
using a different
gas composition from the ones used to introduce dissolved nitrogen or the same
gas
composition during the entire treatment cycle.
[0009] Also provided are methods for preventing sintering of powders during
the process and
promoting nitrogen uptake by the precursor powders. The methods as provided
herein include
continuously agitating the powder to prevent necking or joining between
powders and thus
maintaining supply route of nitrogen around the surface of the powder. The
methods optionally
include providing a rotary hot tube comprising baffles that prevent formation
stratified layers
and continuously breakdown any lumps formed. In other aspects, the methods
include a
fluidized bed reactor that uses the nitrogen containing gas and agitates,
optionally continuously,
the powder mass until the powder is quenched.
[0010] The above and other objects, features and advantages of the present
disclosure will
become more fully understood from the detailed description given herein below
and the
accompanying drawings which are given by way of illustration only, and thus
are not to be
considered as limiting.
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BRIEF DESCRIPTION OF THE DRAWINGS
100111 Exemplary aspects will become more fully understood from the detailed
description
and the accompanying drawings, wherein:
[0012] FIG. lA is a schematic description of the solidification process of
steel involving liquid
to 5-ferrite transformation, followed by austenite and the associated
rejection of nitrogen gas
forming pores;
[0013] FIG. 1B is a schematic description of solidification process of steel
involving liquid to
austenite transformation and the associated retention of dissolved nitrogen
gas in the solid
precursor material according to the teachings of the current disclosure
(exemplary aspect);
[0014] FIG. 2 is a schematic view of an exemplary microstructure of high
nitrogen steel
powder, having nitride precipitates;
100151 FIG. 3 is an exemplary time-temperature cycle for fabricating HNS
powder according
to the some aspects of the current disclosure;
[0016] FIG. 4 is an exempla*, outline of the inventive steps for fabricating
HNS powder
according to the teachings of the current disclosure;
[0017] FIG. 5 is a schematic composition map for adjusting the phase content
in the HNS
powder, according to the teachings of the current disclosure;
[0018] FIG. 6A is a schematic perspective view of an exemplary batch
processing system for
HNS powder wherein the processing chamber includes a rotary tube according to
the teachings
of the current disclosure;
[0019] FIG. 6B is a schematic cross sectional view of an exemplary batch
processing system
for HNS powder wherein the precursor powder is being loaded into the rotary
tube by a screw
feeder according to the teachings of the current disclosure;
[0020] FIG. 7A is a schematic cross sectional view of an exemplary batch
processing system
for HNS powder wherein the precursor powder is being processed within the
rotary tube
according to the teachings of the current disclosure;
100211 FIG. 7B is a schematic cross sectional view of an exemplary batch
processing system
for HNS powder wherein the powder is being removed from the tube after
nitrogen dissolution
and quenched into a collection chamber according to the teachings of the
current disclosure;
[0022] FIG. 8A is a schematic perspective view of an exemplary continuous
processing system
for HNS powder wherein the system includes a heated rotary tube and powder
precursor is
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loaded at one end and the processed powder is collected at the opposite end
according to the
teachings of the current disclosure;
[0023] FIG. 8B is a schematic cross sectional view of an exemplary continuous
processing
system for HNS powder wherein the system includes a heated rotary tube and the
powder
precursor is loaded at one end and the processed powder is collected at the
opposite end
according to the teachings of the current disclosure;
[0024] FIG. 8C is a schematic cross sectional view of an exemplary continuous
processing
system for HNS powder wherein the system includes a heated rotary tube along
with an auger
and the powder precursor is loaded at one end and the processed powder is
collected at the
opposite end according to the teachings of the current disclosure:
[0025] FIG. 8D is a schematic cross sectional view of an exemplary continuous
processing
system for HNS powder wherein the system includes a heated rotary tube having
a vibratory
device and the powder precursor is loaded at one end and the processed powder
is collected at
the opposite end according to the teachings of the current disclosure;
[0026] FIG. 9 is a schematic perspective view of an exemplary fluidized bed
powder
processing system wherein the processing includes gas phase agitation
according to some
aspects of the teachings of the current disclosure;
[0027] FIG. 10 is a schematic cross sectional view of an exemplary fluidized
bed powder
processing system showing the precursor powder being loaded according to some
aspects of
the teachings of the current disclosure;
[0028] FIG. 11 is a schematic cross sectional view of an exemplary fluidized
bed powder
processing system showing the precursor powder being processed according to
some aspects
of the current disclosure;
[0029] FIG. 12 is a schematic cross sectional view of an exemplary fluidized
bed powder
processing system showing the processed powder being quenched according to the
teachings
of the current disclosure;
[0030] FIG. 13 is the schematic process map for powder processing showing
different effects
of temperature and holding time on the powder microstructure according to some
aspects of
the current disclosure;
[0031] FIG. 14 is the schematic cooling rates for powder quenching and their
effects on the
powder microstructure according to some aspects of the current disclosure;
[0032] FIG. 15 presents the Scanning Electron Microscope micrograph for
powders processed
under different process conditions according to some aspects of the current
disclosure;
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[0033] FIG. 16 presents the X-Ray diffraction patterns for as received solid
precursor powder.
and the powders processed under different conditions according to some aspects
of the current
disclosure.
DETAILED DESCRIPTION
[0034] Provided are processes forming high nitrogen steel and devices for
performing the
process. In particular, the processes as provided herein are useful for
creating dissolved
nitrogen in a solid steel material, optionally a powder material. The
processes tailor heating
and optionally holding periods at particular temperatures of the steel to form
various phases,
allow nitrogen to dissolve into the steel and prevent final formation of
nitrides that hinder
corrosion resistance and mechanical strength of the final high nitrogen
steels.
[0035] The following terms or phrases used herein have the exemplary meanings
listed below
in connection with at least one embodiment:
[0036] "HNS" as used herein means steels having high nitrogen content
specifically in
dissolved solid solution form. The amount of nitrogen in the high nitrogen
steel is optionally
equal to or above the amount of nitrogen achievable in an equivalent steel
alloy in a liquid state
at atmospheric pressure of nitrogen.
100371 "Precursor" as used herein means the starting steel powder used to make
the HNS
powder where the precursor powder has a lower nitrogen content that the
resulting HNS
powder.
100381 "Compound" as used herein, means a material formed by reactions between
elements
having a stoichiometric ratio, illustratively, Cr2N and Fe2N, etc.
[0039] "Solid solution" as used herein, means an alloy formed by dissolving
one or more
alloying element(s) in a host solid without changing its phase. In specific
aspects as provided
herein, y-Fe[N], wherein N is the alloying element dissolved in FCC-Fe, the
austenite phase.
[0040] The addition of nitrogen improves the strength, ductility and impact
toughness in
austenitic steels, while the fracture strain and fracture toughness are not
affected at elevated
temperatures. The strength of nitrogen alloyed austenitic steels arises from
three components:
strength of the matrix, grain boundary hardening, and solid solution
hardening. The matrix
strength is not appreciably impacted by nitrogen, rather matrix strength
correlates to the friction
stress of the face centered cubic (FCC) lattice that is mainly controlled by
the solid solution
hardening of the substitutional elements like chromium and manganese. Grain
boundary
hardening, however, which occurs due to dislocation blocking at the grain
boundaries,
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increases proportionally to the alloyed nitrogen content. The highest impact
on the strength
results from the interstitial solid solution of nitrogen. Nitrogen increases
the concentration of
free electrons promoting the covalent component of the interatomic bonding and
the formation
of Cr-N short range order (SRO). The occurrence of Cr-N SRO and the resultant
interactions
with dislocations and stacking faults are believed to play a major role in the
deformation
behavior of these alloys, and can be tailored to enhance the strength,
ductility, and impact
toughness.
[0041] The composition and temperature strongly influence the stacking fault
energy (SFE)
and in turn, the deformation mechanisms and strengthening behavior of
austenitic steels.
Increasing the SFE, causes the active deformation mechanisms to change and is
generally
favored to achieve pure dislocation glide and enhanced toughness.
Specifically, the effect of N
additions on the SFE in Cr and Mn alloyed steels is reported to be non-
monotonic, exhibiting
a minimum SFE at ¨0.4 wt% N. The decrease in SFE at low N content (e.g. less
than 0.4 wt%)
is believed due to the segregation of interstitial N atoms to stacking faults,
however, at higher
N contents (e.g. at or greater than 0.4 wt%) the SFE increases as the bulk
effect of interstitial
solid solution becomes more pronounced. However, the formation of nitrides
such as Cr2N, at
elevated N content, affects the distribution of alloying elements within the
lattice and in turn
diminishes the bulk effect of interstitial solid solution and the SFE. The
formation of nitrides
occurs when the nitrogen content goes beyond certain threshold value (depends
on the overall
composition of the alloy) and should be discouraged to take advantage of the
interstitial solid
solution hardening phenomenon described above.
100421 As such, the high nitrogen steels of the present disclosure are
optionally free or
substantially free of any nitrides. In steels containing alloying elements
(e.g. Cr, Al, Mo, V,
Ti, etc.) nitride formation occurs because these alloying elements are
stronger nitride formers
than iron. As such, nitrides of the type MxNy (where M is Cr, Al, Mo, V. Ti,
etc., x any y are
chosen to arrive at proper stoichiometry) develop with more propensity. The
high nitrogen
steels produced by the processes as provided herein are optionally absent,
optionally
substantially absent a nitride of Cr, Al, Mo, V, Ti, or others.
100431 High nitrogen containing austenitic steels also exhibit excellent
resistance to
atmospheric corrosion. However, the corrosion resistance is also strongly
influenced by the
nitrogen content. At low N content, the formation of a phase (an intermetallic
compound with
Cr) at the grain boundaries as well as the formation of nitrides such as Cr2N
at high nitrogen
content are detrimental to the corrosion resistance of these steels. Best
corrosion resistance can
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be achieved if all nitrogen is in solid solution, i.e. no nitrides such as
Cr2N are precipitated. It
can be summarized that an optimal combination of toughness and corrosion
resistance can be
achieved by limiting the nitrogen content within a range, wherein a
substantially or completely
precipitation free homogeneous microstructure with N in solid solution form
can be obtained.
This range of dissolved N depends on other alloying elements present in the
alloy as well as
the process thermal history as discussed herein.
[0044] One approach to obtain a homogeneous dissolved nitrogen content in a
steel alloy,
specifically in austenitic steel is to (i) dissolve the nitrogen into the
alloy in liquid state and
then (ii) solidify the alloy without losing the dissolved nitrogen during
solidification. However,
both the tasks have their own challenges. For example, the nitrogen solubility
in liquid iron at
atmospheric pressure is very low (0.045 wt% at 1600 C). Nitrogen solubility
in a liquid alloy
increases by the square root of the partial pressure (Sievert's square root
law). Hence, to
introduce higher nitrogen into liquid iron/steel, melting should be done in a
high pressure
nitrogen environment. Nitrogen alloying in the molten state may be achieved by
high pressure
induction or electric arc furnaces, pressure electro slag remelting furnace
(PERS), and plasma
arc and high-pressure melting with hot isostatic processing (HIP), etc.
[0045] Further, it is also known that the addition of certain elements such as
chromium,
manganese vanadium, niobium, and titanium increases the nitrogen solubility,
while addition
of elements such as carbon, silicon, and nickel reduces the nitrogen
solubility. Hence, in order
to induce high nitrogen concentrations into the melt, chromium and manganese
can be added
and nickel should be avoided. Furthermore, in some aspects, elements such as
vanadium,
niobium, and titanium, are absent or present in insignificant amounts as they
are powerful
nitride formers.
[0046] While chromium addition significantly enhances nitrogen solubility in
the melt, it is
also a strong 5-ferrite stabilizer. As illustrated in FIG. 1A, 6-ferrite
solidification in iron alloys
is associated with a wide solubility gap and a sudden drop 12 of nitrogen
solubility in the
material. In other words, a melt containing dissolved nitrogen 13, will lose
most of its nitrogen
during 6-ferrite solidification even though the subsequent lower temperature
austenite phase
can dissolve a much higher amount of nitrogen, 11. It is important to note
that in ferritic steels,
enhancing the dissolved nitrogen content 14 in the liquid by alloying
additions and performing
the melting operation under high nitrogen pressure would not retain the
dissolved nitrogen in
the 8 phase due to the associated loss during 6-ferrite solidification. This
leads to the formation
of interdendritic pores 18, which results in degraded material quality and the
loss of nitrogen
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in the final material. Therefore, to retain the enhanced dissolved nitrogen
achieved through
high nitrogen pressure melting and alloying adjustment and transfer it to the
solid austenitic
material, the 5-ferrite solidification should be avoided. However, if the
solidification operation
is carried out under high nitrogen partial pressure, the pores can be
suppressed increasing the
N content to some extent 15, and importantly after the 5¨>7 transformation,
substantial amount
of nitrogen 16 can be dissolved in the 7 phase; the extent of which depends on
the holding
temperature, pressure and time.
[0047] Now referring to FIG. 1B, in the absence of &ferrite solidification,
wherein the liquid
directly solidifies into austenitic material, much of the dissolved nitrogen
12 in the liquid state
will be retained in the mixture of austenite and the liquid 18' and
subsequently transfer into
the solid austenite phase 13'. It is to be noted that the austenite phase can
have a significant
amount of dissolved nitrogen 11' and in order to achieve the saturation level
11' the liquid may
contain higher dissolved nitrogen 14' to start with, which can be achieved
only by high pressure
melting and alloying adjustment. Further, under high nitrogen partial pressure
the austenite can
pick up more nitrogen 16' and depending upon the temperature and time of
holding the nitrogen
content can reach the solubility limit 11'. The elimination of 5-ferrite
solidification step can be
achieved by carefully adjusting the composition of the alloy. To this end,
manganese addition
plays an important role. While enhancing the nitrogen solubility in the melt,
manganese also
suppresses the formation of 5-ferrite during solidification. As discussed
above, the significant
enhancement of strength in nitrogen alloyed austenitic steel comes from the
formation of Cr-
N SRO. Additionally, Cr enhances the resistance against atmospheric corrosion
and hence is
an important alloying addition. Further, the effect of manganese on enhancing
nitrogen
solubility is known to be two times less than the effect of chromium. Hence,
significantly
higher amount of Mn compared to Cr may be present in order to provide
equivalent nitrogen
solubility, eliminate &ferrite formation as well as achieve enhanced toughness
and corrosion
resistance. Another way to promote austenitic solidification and avoid
degassing of nitrogen is
to add carbon, however, carbon contents > 0.1 wt.% have negative influence on
corrosion
resistance and ductility of the material and hence should be avoided.
[0048] One main problem for the production of austenitic steels containing
high manganese is
the strong segregation behavior of manganese that leads to heterogenic
microstructure; which
is detrimental to the mechanical behavior as well as corrosion resistance.
Further, as discussed
above, precipitation of a phase or nitrides such as Cr2N 24 as shown in FIG. 2
should be
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avoided during processing to achieve high toughness and corrosion resistance.
The segregation
and precipitation issues can be suppressed or completely eliminated by rapidly
cooling.
[0049] As a way of background, when austenitic steel is exposed to a nitrogen
atmosphere at
high temperature, nitrogen may be incorporated in steel through dissolution in
the austenitic
phase up to its solubility limit, according to equation (1):
14 N2(gas) [Nly (1)
and through precipitation of chromium nitrides, according to equations (2) and
(3):
[City + [Nly a [CrNly (2)
2[Cr]y + [N]y Cr2N (3)
Nitrogen in Eqn (1) remains in solid solution depending on temperature of
thermo-chemical
treatment and nitrogen pressure. Nitrogen loss or nitrogen pickup may occur
according to
Sieverts' law at a given set of temperature and nitrogen partial pressure
parameters. Also, note
that CrN in Eqn. (2) is a coherent precipitate and is beneficial in enhancing
mechanical
properties of the steel, whereas Cr2N is a precipitate that deteriorates the
corrosion resistance.
Further, due to slow diffusion rate of N in steel, nitrogen pick up is fast
when the surface area
is large which is achieved by exposing the powder to a nitrogen atmosphere.
However, the
presence of oxide layer on the steel surface inhibits the diffusion of
nitrogen and should be
removed for enhance the rate of diffusion. This can be achieved by treating
the powder in a
reducing gas atmosphere. Once the nitrogen is dissolved in the steel, it can
be retained in the
powder by quenching the powder to a temperature where the nitrogen diffusion
is virtually
absent.
[0050] Provided are methods for making nitrogen steel powder with a dissolved
nitrogen
content, the said dissolved nitrogen content optionally higher than the
solubility limit of N in
the alloy in its liquid state at atmospheric pressure and optionally the
nitrogen alloy powder
being devoid of a nitride compound precipitates or nitride compound layer.
Referring to FIGS.
3 and 4, exemplary methods for the fabrication of the nitrogen alloy powder
are provided. The
temperature and time cycles 30 and method 40 may include one or more of the
following steps;
providing a solid precursor, optionally in powder form, surface form, coating
form or other,
the solid precursor steel with a substantially low dissolved nitrogen content
in step 41 and the
disposing the solid precursor powder into a nitrogen gas or mixture of a
reducing gas and
nitrogen gas environment in steps 42-44, where the said precursor powder can
undergo an
exemplary temperature-time cycle 32 or 33 to obtain the nitrogen steel powder
with a dissolved
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nitrogen content in step 45. Further descriptions on steps 42-44 and
temperature-time cycles
32 and 33 are provided below.
10051] The reducing gas can optionally be a mixture of nitrogen and hydrogen,
argon and
hydrogen, or anhydrous ammonia. Under the reducing gas, the oxides layers will
be removed
and facilitate nitrogen introduction.
100521 In some aspects, a solid precursor material is in the form of a coating
on another
substrate or other steel type. A coating optionally has a thickness,
optionally from 10 nm to
100 micrometers.
[0053] in some aspects a precursor steel is in the form of a powder. The solid
precursor powder
material can optionally be obtained by atomizing a liquid steel alloy in
atmospheric pressures.
[0054] Optionally, the powder is continuously agitated to provide contact with
the gas as well
as prevent sintering. Various methods for providing continuous agitation are
described in this
disclosure.
100551 The precursor powder has a powder size. The precursor powder size is
optionally
between 5 and 250 micrometers (gm), is optionally between 5 p.m and 150 p.m,
optionally
between 10 p.m and 75 p.m. Powder size is defined as the size that is
appropriately sieved
through a desired sieve where powder below a certain size will pass through a
first sieve and
will have size that will be retained by a smaller second sieve. Choice of
sieve size depends on
the desired powder size.
[0056] The precursor steel is predominantly Fe (i.e. 50 wt% or greater Fe) and
optionally
includes one or more other elements that will promote FCC structure. For
example, a precursor
optionally includes Mn. Mn, when present, may be provided at a weight percent
of 0 to 35.
Optionally, the weight percent of Mn is less than 30. Optionally, the weight
percent of Mn is
19-27. Optionally, the weight percent of Mn is 20-26. The presence of N in
such alloys serves
to promote and stabilize a desired FCC structure even when the amount of Mn or
other FCC
promoting metal is less than 20 weight percent. As such, the dissolved N and
Mn optionally
work in concert to promote austenitic structure to the protective layer metal
alloy. Optionally,
the precursor powder includes Ni, which also promotes austenitic structure.
Ni, when present,
may be provided at a weight percent of 0 to 20. Since Ni reduces the N
solubility in the
protective layer, the Ni is optionally between 0 wt% to 5 wt%. The precursor
powder may
optionally include C, that when present, may be provided at a weight percent
of 0 to 0.2. While
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C improves N solubility, it also reduces the toughness of the resulting alloy.
Optionally, the C
is present in the precursor powder at 0 wt% to 0.1 wt%.
[0057] As mentioned earlier, the strengthening mechanism in nitrogen alloy
steel emerges
from the formation of Cr-N SRO and hence Cr is optionally included in the
provided N alloy.
However, Cr is a 8-ferrite promoter as well as ferrite stabilizer. In order to
control the phase of
the steel, the ferrite stabilizing effect of Cr may be countered by adjusting
the amount of N
and/or Mn, both of which serve as austenite stabilizers. The precursor may
include one or more
other metals. Optionally, a precursor may include molybdenum. Mo, when
present, may be
provided at a weight percent of 0 to 5. Optionally, a precursor may include
aluminum. When
present Al may be provided at 0.01 wt% to 10 wt%. Al is optionally present at
or less than 10
wt%, optionally at or less than 8 wt%, optionally at or less than 6 wt%.
[0058] Now referring to FIG. 3, TN 34d represents a temperature where nitrogen
uptake in steel
occurs through nitride formation. The exact temperature and form of nitride
depends on the
steel composition. Ty 34c represents the temperature at which the steel
transforms into
austenite of FCC form. Again, the temperature depends on the steel
composition. TIN 346
represents the temperature at which all the nitride compounds dissolve and
nitrogen in the steel
exists in dissolved nitrogen form. Tm 34a represents the melting point of the
steel. When the
precursor in step 41 is subject to a temperature-time cycle 32 in steps 42-44,
depending on the
powder size and composition, first nitrogen uptake is expected to occur
through nitride
formation during the ramp-up phase 32', which will dissolve during holding the
powder above
34h. Above 34b, all the nitrogen uptake is expected to occur elemental N form.
The extent of
N uptake will depend on the ramp-up time plus the holding time above 34b. To
retain the N in
the steel in dissolved form and prevent reformation of nitrides, the steel
needs to be quenched
32" below 34d, quickly, which can be achieved by various methods. The
quenching rate
influences the final microstructure of the HNS. Under slow cooling the N may
precipitate into
carbide particles. The gas pressure, temperature and time are adjusted
according to the desired
dissolved nitrogen content in the final HNS of step 45 and the composition of
the precursor.
[0059] Introducing dissolved N directly into a given steel powder without the
formation of
substantial nitrides following temperature-time cycle 32 (FIG. 3) may present
some challenges.
The desired holding temperature in this case is too close to the melting
temperature, Tm 34a of
the alloy. When precursors are disposed to this temperature range, they tend
to join together by
a process known as sintering in the art. Although according to this disclosure
the presence of
sintering is not always undesirable, in many aspects sintering may hinder the
use of any
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resulting powder in a coating application process, and hence undesirable in
some applications.
Further, sintering hinders compositional homogeneity of the powders,
particularly in regards
to N. Holding the temperature in this range for a short period without
sintering, however, would
limit the uptake of N. In other words, for a low target N uptake one can adopt
temperature-time
cycle 32 without any detrimental effects due to sintering. However, to avoid
sintering but to
achieve relatively increased nitrogen uptake, optionally, one can adopt
alternative temperature-
time process.
100601 Referring again to FIG. 3, following temperature-time cycle 33, one can
optionally hold
the powder in the nitrogen environment just above the nitride formation
temperature, TN for a
nitrogen uptake time to facilitate the uptake of significant amount of N.
Optionally, a
temperature for a nitrogen uptake time is at or above the TN and below the Tr
of the alloy.
Further, holding the powder at this temperature prevents sintering especially
in the presence of
powder agitation is used during the nitrogen uptake time as will be
illustrated below. However,
holding the powder at a temperature between the Ty and TN may lead to the
formation of
undesirable nitrides. Accordingly, according to some aspects, the temperature-
time cycle 33,
comprises of a second heating step, where the temperature is raised above TyN
and held there
for a brief nitride conversion time to decompose the nitrides, yielding the
desired dissolved
nitrogen content, prior to quenching the powder. The nitride conversion time
above T7N is
substantially shorter compared to nitrogen uptake time and thus prevents
sintering. It will be
appreciated that TN, Ty, TN depend on the steel composition, and further the
holding times also
depend on the powder composition and size. The temperature-time cycles 32 and
33 are
exemplary illustrations and many variations can optionally be adopted to
achieve sinter/nitride
free powder with a desired dissolved N content.
100611 In some aspects, a nitrogen uptake time is in excess of I second.
Optionally, a nitrogen
uptake time may be indefinite, but is more commonly 1 hour or less. For larger
steel pieces or
larger powder sizes the nitrogen uptake time may be adjusted upward. In
particular aspects, a
nitrogen uptake time is from 1 second to 15 minutes, optionally 1 second to
100 seconds,
optionally 1 second to 60 seconds, optionally 10 seconds to 100 seconds,
optionally 30 seconds
to 70 seconds, optionally 50 seconds to 60 seconds. A hold time may be
sufficient to fully heat
the precursor to the desired temperature or may hold the precursor at that
temperature for the
nitrogen uptake time.
100621 A nitrogen uptake time may be at a constant temperature or may be at a
varying
temperature. A varying temperature during a nitrogen uptake time may be at or
between TN
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and T. The temperature may fluctuate or remain substantially constant,
optionally varying by
C or less.
[0063] The precursor with the nitrogen uptaken into the material may then be
subjected to a
further heating step whereby the precursor is heated to a temperature near,
but not at or above
the Tm. Optionally, the second heating step heats the precursor to a second
temperature that is
above a TN of the precursor and below a melting temperature for the precursor
powder. As
illustrated above, above the TN of the precursor, any nitrides formed in the
uptake step or
otherwise present in the steel material are converted into dissolved nitrogen.
This further
increases the weight percent of dissolved nitrogen and prevents unwanted
characteristics that
occur due to the presence of nitrides in the final high nitrogen steel.
[0064] The increased temperature above the TIN of the precursor is optionally
held for a nitride
conversion time. A nitride conversion time is optionally any time to allow
all, substantially all
or any desired amount of nitrides within the precursor to be converted to
dissolved nitrogen. A
nitride conversion time is optionally 1 hour or less. In some aspects, a
nitrogen conversion time
is a short as possible so as to both convert the nitrides to dissolved
nitrogen but also to prevent
sintering (in some aspects). As such, a nitride conversion time is optionally
less than 1 hour,
optionally less than 20 minutes, optionally, less than 10 minutes, optionally
less than 5, 4, 3, 2,
1, min. It has been observed that some sintering may occur when using
particular precursor
steel in the form of a powder at 10 minutes. As such, when a powder precursor
is used, the
nitride conversion time is optionally 10 minutes or less, according to some
non-limiting
aspects.
[0065] Once the desired dissolved nitrogen content is achieved in step 43, the
powder is
quenched to a temperature where the diffusion is virtually absent in step 44
following an
exemplary temperature-time cycle 33. The cooling rate is critical to avoid
nitride formation or
reformation. Optionally, the powder cooling rate is between 1 Cis and 100
'Cis, optionally
the powder cooling rate is between 5 'Cis and 50 'Cis, optionally the powder
cooling rate is
above 10 C/s.
[0066] The atmospheric pressures used in the processes optionally are not
required to exceed
1 atm as, in many aspects, 1 atm is sufficient to radically increase the
amount of dissolved
nitrogen in the high nitrogen steel relative to prior processes. However, in
other aspects the
atmospheric pressure is optionally above 1 atm, optional 2 atm or greater,
optionally 3 atm or
greater, optionally 4 atm or greater. In some aspects, the atmospheric
pressure is less than that
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typically required to dissolve nitrogen in liquid steel. As such, an
atmospheric pressure is
optionally less than 10 atm, optionally less than 9 atm, optionally less than
8 atm, optionally
less than 7 atm, optionally less than 6 atm, optionally less than 5 atm.
100671 The resulting high nitrogen steel is provided with an exceptionally low
nitride content,
optionally 0.01 wt% or lower. Optionally, the resulting HNS has a nitride
content at or below
0.02 wt%, optionally 0.03 wt%, optionally 0.04 wt%, optionally 0.05 wt%,
optionally 0.1 wt%
or lower.
[0068] The resulting high nitrogen steel optionally has a dissolved nitrogen
content of 0.05
wt% to 6.0 wt%, or higher or any value or range therebetween. Optionally the
dissolved
nitrogen content of the HNS is at ore greater than 0.1 wt%, optionally 0.5
wt%, optionally 1
wt%, optionally 2 wt%, optionally 3 wt%, optionally 4 wt%, optionally 5 wt%,
optionally 6
wt%. In many aspects the amount of dissolved nitrogen exceeds the solubility
limit of nitrogen
in the alloy (alloy of otherwise identical composition) in a liquid state at
atmospheric pressure.
[0069] In some aspects, the resulting high nitrogen steel includes a ferrite
(a) phase, austenite
(7) phase, or a mixture of a + 7 phase. In some aspects, the alloy is
predominantly a single
phase. A single phase may be a ferrite phase or a gamma phase. Optionally, an
alloy is
predominantly or entirely a single phase, optionally a 7 phase. The HNS
produced by the
processes as provided herein is optionally predominantly FCC structure,
optionally 90% or
greater FCC structure. Optionally, the HNS produced by the processes as
provided herein are
absent BCC structure throughout the HNS.
100701 As discussed above, some elements act as austenite stabilizers while
others promote
ferrite. Further, the extent of their influence also varies considerably. For
example, N is almost
20 times more effective in stabilizing austenite compared to Mn. Similarly, Cr
is almost two
times more effective than Mo in stabilizing ferrite. Therefore, to predict the
phases of the iron
alloys of this disclosure, it is appropriate to use a nitrogen equivalent as a
predictor of
austenite/ferrite composition in a N alloyed protective layer as presented in
this disclosure. For
iron alloys primarily containing Mn, Cr, and N alloying elements, the N and Cr
equivalents can
be expressed as: N_eq = 10 (wt.% N) + 0.25 (wt.% Mn) ¨ 0.02(wt.% Mn)2 +
0.00035(wt.%
Mn)3 and Cr_eq = wt.% Cr, respectively. Note that should any other elements be
present in
appreciable amount, whether austenite stabilizer or ferrite stabilizer, N_eq
and Cr_eq is
modified appropriately. Further, there is a lot of controversy regarding
weight factors for each
element and often they are empirically determined from experiments. But, there
is a general
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agreement that N and C are the two most impactful austenite stabilizers. Since
addition of C
beyond 0.1 wt% is detrimental to the toughness, primarily the influence of N
and Mn is
considered here for exemplary illustration of alloy compositions.
100711 Accordingly, the target alloy composition impact on phase stability is
illustrated in FIG.
5, wherein the phase boundary between 100% austenite and the mixture austenite
+ ferrite is
separated by a line which can be expressed as N equivalent = A x Cr equivalent
- B. Based on
experimentations, A is ¨0.98 and B is ¨11.5 for Ni free Fe-Mn-Cr-N alloy.
Accordingly,
exemplary alloy compositions will lead to the following outcomes as presented
in Table 1. The
impact of Mn content in stabilizing the austenite decreases as the content
increases. For
example, keeping the nitrogen concentration at 0.5 wt%, an increment of Mn
content from 15
wt% to 30 wt%, decreases the N-eq from 5.27 to 3.65. Further, N concentration
is the most
influential factor in stabilizing the austenite. For example, by changing the
N concentration
from 0.5 wt% in alloy #4 to 0.7 wt% in alloy #5, results in an austenitic
alloy even though
significant amount of Cr (20 wt%) is present in the alloy. However, care must
be taken not to
increase the N content significantly beyond the stability zone especially when
high amount of
Cr is present to prevent Cr2N precipitation as illustrated in FIG. 2.
Alternatively, Mn addition
can counter the influence of Cr and contribute towards the stability of
austenite. Optionally,
the N kept between 0.4 wt.% and 0.9 wt.%, Mn is kept between 19-27 wt% and the
Cr is kept
between 10-18 wt.%, the rest being iron.
Table 1
Alloy # N(wt%) Mn(wt%) Cr(wt%) =N eq(wt%) Cr eq(wt 10) Phase
1 0.5 15 13 5.27 13
2 0.5 20 13 4.6 13
3 0.5 30 13 3.65 13
4 0.5 20 20 4.6 20
0.7 20 20 6.6 20
[0072] An exemplary alloy containing 15 wt% Cr, 25 wt% Mn and 0.7 wt% N and
the
remainder Fe would form an austenite phase which is preferred in many
applications. In some
aspects, a N alloy is or includes 13-14 wt% Cr, 20-26 wt% Mn, and 0.4-0.6 wt%
N with the
remainder being Fe.
[0073] Accordingly, the exemplary method described in FIGS. 3 and 4, can be
practiced in
various embodiments as illustrated in FIGs. 6-12 as are described herein
below.
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100741 According to the teaching of this disclosure, an exemplary embodiment
60 for batch
processing of powder shown in FIGS. 6A and 6B, includes a ceramic processing
tube 65 which
is rotated using a motor-gear arrangement 67. One end of tube 65 is connected
to a powder
feeding apparatus, which comprises of motor 61, precursor inlet 62 to load
powder and an auger
63. This feeding apparatus can be inserted or retracted using lead screw
arrangement 69 to
evenly distribute powder through the processing tube 65. The other end of the
tube 65 is
operably coupled to a cooled collection chamber 66, which can be cooled by
chilled gas or
other fluid. The powder treatment tube is heated to a desired temperature by a
heater 64.
100751 Further details of processing and removal of the precursor is shown in
Fig. 7A and Fig.
7B respectively. To enhance processing of the precursor powder, fins 79 can be
added to the
processing tube 75. In order to prevent any oxidation of the powder, the
internal environment
of the processing tube is always kept under nitrogen atmosphere. The fins can
be ceramic rods
placed along wall of the processing tube 75 as shown in cross section 78. In
order to remove
the precursor, a tilt angle is applied to the rotating processing tube 75
through a jack
arrangement 72 which in turn empties the treated powder into the collection
chamber 76.
100761 Another exemplary embodiment 80 for continuous processing of powder
shown in
FIGs. 8A and 8B, includes a ceramic processing tube 85 which is rotated using
a motor-gear
arrangement 87. One end of tube 85 is operably connected to a powder feeding
apparatus 81,
via precursor inlet 82 to continuously feed powder. The other end of the tube
85 is operably
coupled to a cooled collection chamber 86, which can be cooled by chilled gas
or other cooling
fluids. The powder treatment tube is heated to a desired temperature by a
heater 84. The tube
assembly can be placed at different angles relative to gravity via an
adjustment device 88. This
angle as well as the rotation speed of the tube 85 determines the residence
time of the precursor
powder in the treatment tube. Accordingly, various temperature-time cycles
presented in FIG.
3 can be adopted through this arrangement to continuously treat the powder to
achieve a desired
dissolved N content in the powder.
100771 In another exemplary embodiment 80', the auger 87' extends from the
powder reservoir
82' till the delivery end of the processing tube 85'. Optionally, the
processing tube remains
stationary, while the auger continuously agitates the powder inside the
processing tube.
Optionally, the processing tube also rotates while treating the powder. The
pitch and rotational
speed of the auger 87' controls the feed rate of the precursor as well as the
dwell time in the
processing tube 85'. The auger eventually pushes the precursor into the
collection chamber 86'.
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In such a system, auger is made of high temperature compatible materials such
as ceramic to
be able to operate at high temperatures in the processing tube.
[0078] In yet another exemplary embodiment 80", precursor feeding can be
achieved using
electromagnetic vibration 81". The feed rate of the precursor is controlled by
regulating the
vibration frequency. The feed rate can be further increased by tilting the
setup using a jack
arrangement 83". The precursor is introduced into the processing tube 85"
using commercially
available powder feeder 82" such as thermal spray powder feeder. Due to
continuous vibration
of the processing tube 85", the precursor powders are agitated and moved
through the
processing tube into the collection chamber 86".
[0079] Referring to FIG. 9. embodiment 90 provides a methodology to alloy the
said precursor
with N via a fluidized bed reactor. Embodiment 90 includes a heating element
92, a processing
tube 93, a precursor intake 91, a vent for the exhaust gas 95, a downstream
gate valve 98 and
a cooled collection chamber 97, operably connected to the processing tube 93
via gate valve
98. As illustrated in the cross sectional view FIG. 10, a known quantity of
precursor 99 is
introduced into the processing tube 93 via the intake port 91, while keeping
the gate valve 98
closed. Subsequently, preheated nitrogen gas is continuously injected through
nozzle 94 which
fluidizes the precursor as shown in FIG. 11. The exhaust gas leaves the
processing chamber
through the vent 95, which can be recycled back to the chamber nozzle 94. The
combined heat
input from the heating element 92 and the preheated fluidizing gas keeps the
precursor powder
temperature at a desired range. The temperature-time cycle to achieve a target
dissolved
nitrogen content in the steel powder is adopted according to the teachings of
this disclosure as
illustrated in embodiments 30 and 40. After the desired N concentration is
achieved in the
powder, as shown in FIG. 12, the downstream valve 98 is opened to release the
treated powder
99 into the collection chamber 97 through the cooled channel 96. It is
understood that the
fluidized bed treatment method as described herein can be achieved by
alternative
embodiments following similar mechanical and physical principles.
Example
[0080] A precursor powder with composition of Fe, 12 wt% Cr and 20 wt% Mn was
centrifugally atomized at Ervin Technologies, Tecumseh, Michigan, USA and
classified to
yield a size distribution of 10 p.m to 60 pm. The powder was then processed
according to the
teachings of the present disclosure by utilizing an embodiment illustrated
FIG. 8C. Nitrogen
uptake in the powder was determined using Leco TC436 combustion analyzer at
NSL
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Analytical Services, Inc., Ohio, USA. The phase of the powder was determined
using a Rigaku
Mini flex X-ray Diffractometer (Cu Ka radiation ?l .5402A).
[0081] The precursor powder was subject to various temperature-time cycles
under N2+5%H2
gas mixture environment. The observations are presented in FIGS. 13-16. As
illustrated in FIG.
13, when the process temperature remained between Ti (<980 C) and T7, for
example 930 C,
nitride precipitated were always present in the final powder as shown FIG.
15D. In the early
stages of the treatment process, especially when the treatment temperature was
closer to Ty,
both a and 7 phases along with grain boundary precipitates were observed as
shown in FIG.
15B. The corresponding x-ray diffraction pattern B shown in FIG. 16, indicates
the presence
of a andy phases. It is to be noted that the incoming precursor powder had
single phase (a or
ferrite), as demonstrated by the diffraction pattern A in FIG. 16. With
increasing treatment
time, the overall nitride precipitate content increased and the matrix alloy
transformation into
7 phase. Longer holding time led to powder sintering. Processing temperatures,
between Ty and
TN, formed nitride precipitates in the a matrix. No changes were observed
below the threshold
temperature Tth. When the process temperature remained above T2 (>1080 C), for
example,
1130 C, although the N uptake was accelerated, the powders sintered quickly.
The
corresponding powder microstructure is shown in FIG. 15A. Temperature¨time
cycle within
the shaded region as illustrated in FIG. 13, always resulted in nitride free
single phase powder
as shown in FIG. 15C and the corresponding X-ray diffraction pattern is shown
FIG. 16A. It is
to be understood that the boundaries of this operational regime is not
limiting and is based on
set of experiments, which is not completely exhaustive. It is to be noted that
factors such as
alloy composition, powder size and agitation during the treatment would affect
the area of this
zone. Powder cooling rate played a very important role as illustrated in FIG.
14. Cooling rate
of 1 C/s always led to precipitate formation although the treated powder
started with a single
phase homogenous alloy. Cooling rate of 10 Cls ensured precipitate free single
phase powder
at room temperature, provided the treated powder started with a single phase
homogenous
alloy. Accordingly, a treatment zone as illustrated in FIG. 13, can be
constructed to achieve
single phase homogeneous dissolved nitrogen alloy according to the teachings
of this
disclosure. As shown in Table 2, processing at 980 C and 1080 C for a dwell
time of 300s,
followed by quenching to room temperature at 10 C/s, yielded nitrogen content
of 0.3 wt% and
0.48 wt% respectively, without any precipitates and uniform austenite phase.
Further, sintering
of powder was also absent in these experiments.
Table 2
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Temperature C Processing time (s) Nitrogen uptake %wt Quench rate 'Cis Phase
980 300 0.3 10
1080 300 0.48 10
[0082] While
particular embodiments have been illustrated and described herein, it
should be understood that various other changes and modifications may be made
without
departing from the scope of the claimed subject matter. Moreover, although
various aspects of
the claimed subject matter have been described herein, such aspects need not
be utilized in
combination. It is therefore intended that the appended claims cover all such
changes and
modifications that are within the scope of the claimed subject matter.
[0083] Various
modifications of the present invention, in addition to those shown and
described herein, will be apparent to those skilled in the art of the above
description. Such
modifications are also intended to fall within the scope of the appended
claims.
[0084] It is
appreciated that all reagents are obtainable by sources known in the art
unless otherwise specified.
[0085] The
foregoing description is illustrative of particular embodiments of the
invention, but is not meant to be a limitation upon the practice thereof.
[0086] Various modes for carrying out the present invention are disclosed
herein; however, it
is to be understood that the disclosed modes are merely exempla*, of the
invention that may
be embodied in various and alternative forms. The figures are not necessarily
to scale; some
features may be exaggerated or minimized to show details of particular
components.
Therefore, specific structural and functional details disclosed herein are not
to be interpreted
as limiting, but merely as a representative basis for teaching one skilled in
the art to variously
employ the present invention.
[0087] Reference is made in detail to compositions, aspects and methods of the
present
disclosure. It is also to be understood that this disclosure is not limited to
the specific aspects
and methods described herein, as specific components and/or conditions may, of
course, vary.
Furthermore, the terminology used herein is used only for the purpose of
describing particular
aspects of the present disclosure and is not intended to be limiting in any
way.
[0088] it must also be noted that, as used in the specification and the
appended claims, the
singular form "a," "an," and "the" comprise plural referents unless the
context clearly indicates
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otherwise. For example, reference to a component in the singular is intended
to comprise a
plurality of components unless explicitly noted otherwise.
[0089]
Throughout this description, where publications are referenced, the
disclosures
of these publications in their entireties are hereby incorporated by reference
to more fully
describe the state of the art to which this disclosure pertains.
21
