Note: Descriptions are shown in the official language in which they were submitted.
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PROCESSING OF INTERMETALLIC ALLOYS
BACKGROUND OF THE INVENTION
Field of the Invention:
The present invention is directed to processing of intermetallic alloys such
as aluminide
alloys. More specifically, the present invention is directed to the processing
of iron
aluminides by powder metallurgical techniques. These techniques result in
processing routes
optimized to take advantage of the dynamic recrystallization or superplastic
behavior of the
alloys .
State. of the Art:
In the discussion of the state of the art that follows, reference is made to
certain
structures and/or methods. However, the following references should not be
construed as an
admission that these structures and/or methods constitute prior art. Applicant
expressly
reserves the right to demonstrate that such structures and/or methods do not
qualify as prior
art against the present invention.
Ordered intermetallic alloys based on iron aluminide Fe3Al offer a combination
of
attractive properties such as excellent resistance to oxidation and
sulfidation at high
temperature, and high strength to weight ratio. (See, for example, R. S.
Sundar, et al.,
Mater. Sci. and Eng. A, 258, (1998) 219-228.) They are also potential low-cost
replacements
for more expensive high temperature structural alloys containing strategic
elements such as
nickel and chromium. Although iron aluminide alloys exhibit limited room
temperature
ductility, poor high temperature strength, low fracture toughness, poor
machinability and poor
resistance to environmental embrittlement, substantial improvement in these
properties can be
achieved by a combination of composition and process control.
More recently, intermetallics based on nickel, iron and titanium aluminides
have been
the subject of research due to their excellent thermal stability at high
temperatures coupled
with their unique combination of properties such as low densities and good
room temperature
and high-temperature tensile strengths. (See M. R. Hajaligol et al., Mater.
Sci. and Eng. A,
258 (1998) 249-257.) Of the intermetallics, iron aluminides based on FeAI with
a B2
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structure (ordered BCC structure with aluminum atoms occupying the body
centers) are of
more interest than Fe3Al-based alloys. Alloying these alloys with Mo, Zr and
others results in
a combination of attractive properties such as oxidation, corrosion and
sulfidation resistance at
high temperatures. Additionally, the alloy possesses reasonable strengtli at
high temperatures
for use as a structural material. The room temperature ductility of FeAI
alloys are generally
in the range of 2-6 % , and the elongations are influenced by room temperature
embrittlement.
The low ductility of FeAI alloys necessitates hot working of cast materials at
high
temperatures, and hot working approaches limit the manufacturability of sheets
and rods.
Metallurgical processing techniques based on melting and casting, forging and
rolling
have been used to understand the processability of FeAI alloys. Casting
techniques such as
Air Induction Melting (AIM) and Vacuum Induction melting (VIM) have generally
proven
unsatisfactory due to the high porosity of ingots processed by these
techniques. However,
AIM and VIM ingots subjected to a secondary Vacuum Arc Melting (VAM) process
have
exhibited improved porosity and reduced defects. (See C. Testani, et al.,
Proc. Inter.
Symposium on Nickel and Iron Aluminides: Processing, Properties, and
Applications, Eds. S.
C. Deevi, P. J. Maziasz, V. K. Sikka, and R. W. Cahn, ASM International,
Materials Park,
OH, 1996, pp. 213.)
Ingots subjected to hot working display properties that are temperature and
strain rate
dependent, thereby necessitating processing at strain rates lower than
conunercially viable to
avoid unwanted properties such as work hardening. This may be partially
overcome by
extrusion in which low strain rates, in combination with preheated dies and
hydrostatic
pressure, combine to avoid brittle fracture and yet control detrimental grain
growth.
However, all of the above processes suffer from complexity and cost.
It has been shown that the yield strengths, ultimate tensile strengths and
tensile
elongations of hot extruded rods of FeAI alloys are superior to the properties
of the cast FeAI
alloys due to the fine grain microstructure of the hot extruded FeAI alloys.
(See P. J.
Mazsiaz, et al., Proc. Inter. Symposium on Nickel and Iron Aluminides:
Processing,
Properties, and Applications, Eds. S. C. Deevi, P. J. Maziasz, V. K. Sikka,
and R. W. Cahn,
ASM International, Materials Park, OH, 1996, pp. 157.) In addition, it has
been shown that
reactive hot extrusion of Fe and A1 powders improves the properties of fine
grained Fe-24
wt. % A1 over the properties of cast materials. (See S. C. Deevi, et al.,
Proc. Inter.
Symposium on Nickel and Iron Aluminides: Processing, Properties, and
Applications, Eds. S.
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C. Deevi, P. J. Maziasz, V. K. Sikka, and R. W. Cahn, ASM International,
Materials Park,
OH, 1996, pp. 283.) For commercial viability, powder metallurgy is promising
since finer
microstructures with a uniform distribution of micro-constituents can
theoretically be
produced. These, by analogy to other fine grained FeAI alloys, should also
exhibit improved
properties. However, the critical processing paths to achieving these alloys
has not been
clearly delineated.
In view of the above, it is desirable to obtain a simple and cost effective
processing
route to produce iron aluminides with fine grain microstructures and which
ensures high
quality defect-free products on a repeatable basis.
Due to the irregular shapes of water-atomized powders, they can be cold
compacted
and continuously rolled to produce good quality sheet products. In contrast,
gas atomized
powders have spherical shapes and require hot compacting methods to
consolidate them. Hot
extrusion of these powders has resulted in semi-products with fine-grained
microstructures
with attractive properties. For developing components from these semi-
products, it is
important to characterize their hot deformation behavior so as to help in
designing metal
working processes.
Different approaches available for evaluating the hot working mechanisms
include
evaluation of the shapes of the stress-strain curves, the standard kinetic
parameters and the
processing maps. The kinetic analysis of hot deformation uses the rate
equation:
E =A6~exp[-Q/RT] (1)
where is the strain rate, A is a constant, 6 is ~
the flow stress, ya is the stress exponent, Q is the activation energy, R is
the gas constant and T
is the temperature. The kinetic parameters - the stress exponent and the
apparent activation
energy - are evaluated from experimental data and compared with the
corresponding values
known for some atomistic mechanisms for identifying the rate controlling one.
The limitations
of this analysis have been its applicability to complex alloys and its
inability to optimize the
hot workability.
The use of processing maps for optimizing hot workability and controlling
microstructure during mechanical processing has been found to be very
beneficial. (See, for
example, Y. V. R. K. Prasad and S. Sasidhara, Eds., Hot Working Guide: A
Compendium of
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Processing Maps, ASM International, Materials Park, OH, 1996.) With the help
of a
processing map, it is possible not only to arrive at the optimum parameters
for designing a
metalworking process without resorting to expensive and time consuming trial
and error
methods but also to control the grain size.
In brief, depicted in a frame of temperature and strain rate, power
dissipation maps
represent the pattern in which the power is dissipated by the material through
microstructural
changes. The rate of this change is given by a dimensionless parameter called
the efficiency,
r~, of power dissipation:
T~ = 2rnl(m + 1 ) (2,)
where m is the strain rate sensitivity of flow stress. Over this frame is
superimposed a
continuum instability criterion for identifying the regimes of flow
instabilities, developed on
the basis of extremum principles of irreversible thermodynamics as applied to
large plastic
flow and given by another dimensionless parameter:
~(E) _ [aln(mlm +1 )/alnE] +m
wherein is the applied strain rate. Flow
instabilities are predicted to occur when ~ is negative. These two maps
together constitute a
processing map, which exhibits domains with local efficiency maxima
representing certain
specific microstructural mechanisms and also regimes of flow instabilities.
The exploitation of powder metallurgy routes is very promising since a much
finer
microstructure with a uniform distribution of microconstituents can be
produced by this
technique. In addition, the mechanical processing of compacts formed by powder
metallurgical methods is an important first step in view of the limited
workability of the
powder metallurgy compacts.
BRIEF SUMMARY OF THE INVENTION
The present invention overcomes the deficiencies previously associated with
conventional casting and powder metallurgical technologies for the processing
of intermetallic
alloys.
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In addition, the invention provides a method of manufacturing a worked product
from
an intermetallic alloy such as iron, nickel or titanium aluminide alloy which
produces sound
material retaining all of the advantageous material properties commonly
associated with these
alloys while providing the added advantage of lower cost.
More particularly, the present invention provides for a method of manufacture
comprised of preparing a body from an intermetallic alloy powder, preferably
an iron
aluminide, the powder formed by either water atomization or gas atomization
techniques, and
hot working the body at a strain rate of 0.001 to 1.0 s-1 and at a temperature
range above
750 ° C during which the intermetallic alloy undergoes grain
refinement. Alternatively, the
body is forged or rolled at a temperature of 1100 to 1250°C prior to
the hot working step.
The strain rate, the selection of which is aided by the use of power
dissipation maps, is
sufficient to achieve either dynamic recrystallization or superplastic
deformation of the
intermetallic alloy.
BRIEF DESCRIPTION OF THE DRAWING FIGURES
The objects and advantages of the invention will become apparent from the
following
detailed description of preferred embodiments thereof in connection with the
accompanying
drawings in which like numerals designate like elements and in which:
Figure 1 is a flow diagram illustrating the steps of the method to manufacture
a worked
product from water atomized powders of an intermetallic alloy.
Figure 2 is a flow diagram illustrating the steps of the method to manufacture
a worked
product from gas atomized powders of an intermetallic alloy.
Figures 3 a-b show surface morphologies of (a) water atomized (WA), and (b)
gas
atomized (GA) powders.
Figure 4 shows an initial microstructure of a compacted and extruded billet of
water
atomized Fe-24 weight % A1 alloy.
Figure 5 shows an initial microstructure of a compacted and extruded billet of
gas
atomized Fe-24 weight % A1 alloy.
Figure 6 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 600°C obtained in compression at different strain
rates.
Figure 7 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 650°C obtained in compression at different strain
rates.
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Figure 8 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 700°C obtained in compression at different strain
rates.
Figure 9 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 750°C obtained in compression at different strain
rates.
Figure 10 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 800°C obtained in compression at different strain
rates.
Figure 11 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 850°C obtained in compression at different strain
rates.
Figure 12 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 900°C obtained in compression at different strain
rates.
Figure 13 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 950°C obtained in compression at different strain
rates.
Figure 14 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 1000°C obtained in compression at different strain
rates.
Figure 15 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 1050°C obtained in compression at different strain
rates.
Figure 16 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 1100°C obtained in compression at different strain
rates.
Figure 17 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from WA powder at 1150°C obtained in compression at different strain
rates.
Figure 18 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 850°C obtained in compression at different strain
rates.
Figure 19 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 900°C obtained in compression at different strain
rates.
Figure 20 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 950°C obtained in compression at different strain
rates.
Figure 21 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 1000°C obtained in compression at different strain
rates.
Figure 22 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 1050°C and at different strain rates.
Figure 23 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 1100°C obtained in compression at different strain
rates.
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Figure 24 shows true stress - true plastic strain curves of an iron aluminide
alloy made
from GA powder at 1150°C obtained in compression at different strain
rates.
Figure 25 shows the variation of flow stress with strain rate at different
temperatures at
a strain of 0.5 for a GA FeAI alloy.
Figure 26 is a processing map obtained on an iron aluminide alloy formed from
WA
powders at a strain of 0.1.
Figure 27 is a processing maps obtained on an iron aluminide alloy formed from
WA
powders at a strain of 0.2.
Figure 28 is a processing maps obtained on an iron aluminide alloy formed from
WA
powders at a strain of 0.3.
Figure 29 is a processing maps obtained on an iron aluminide alloy formed from
WA
powders at a strain of 0.4.
Figure 30 is a processing maps obtained on an iron aluminide alloy formed from
WA
powders at a strain of 0.5.
Figures 31 a-d show microstructures obtained on iron aluminide alloy specimens
made
from WA powders deformed at (a) 1100°C/0.001 sn, (b) 1100°C/0.1
sn, (c) 1150°C/0.001 s-
1, and (d) 1150°C/0.1 su.
Figure 32 shows a microstructure obtained on iron aluminide alloy specimen
made
from WA powders deformed at 1150°C and 100 sn strain rate (DRX domain).
Figures 33 a-b show microstructures of iron aluminide alloy specimens made
from WA
powders deformed at (a) 850°C/0.001 s-1, and (b) 900°C/0.1 s-1
Figure 34 relates grain size values measured on iron aluminide alloys made
from WA
powders deformed at different temperatures and strain rates.
Figure 35 is a schematic bifurcation diagram for iron aluminide alloy made
from WA
powders obtained from the changes in the processing maps with strain at a
temperature of
1150°C.
Figures 36 a-c show microstructures obtained on iron aluminide specimens made
from
WA powders deformed at 750°C in the instability regime. (a) 100 s-1,
(b) 10 sn, and (c) 1 s-1.
Figure 37 is an Arrhenius plot for an iron aluminide alloy made from WA
powders in
the domain of dynamic recrystallization.
Figure 38 shows variation of average grain diameter with the Zener-Hollomon
parameter for an iron aluminide alloy made from WA powders.
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Figure 39 is a processing map obtained on an iron aluminide alloy formed from
GA
powders at a strain of 0.1.
Figure 40 is a processing map obtained on an iron aluminide alloy formed from
GA
powders at a strain of 0.2.
Figure 41 is a processing map obtained on an iron aluminide alloy formed from
GA
powders at a strain of 0.3.
Figure 42 is a processing map obtained on an iron aluminide alloy formed from
GA
powders at a strain of 0.4.
Figure 43 is a processing map obtained on an iron aluminide alloy formed from
GA
powders at a strain of 0.5.
Figures 44 a-b show microstructures obtained on iron aluminide alloy specimens
made
from GA powders deformed at a strain rate of 0.1 s-1 and at (a) 1050°C
and (b) 1100°C.
Figure 45 shows variation of average grain diameter of an iron aluminide alloy
made
from GA powders with temperature at a strain rate of O.I s-1. The variation of
efficiency of
power dissipation (at E=0.5) is also shown.
Figures 46 a-b show microstructures of iron aluminide alloy specimens made
from GA
powders deformed at 1100°C and at (a) 0.01 sn and (b) 1.0 s'1
Figure 47 relates grain size values measured on iron aluminide alloys made
from GA
powders deformed at 1100°C at different strain rates.
Figure 48 shows a microstructure obtained on an iron aluminide specimen made
from
GA powders deformed at a strain rate of 100 s-1 and 1050°C,
corresponding to the instability
regime.
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DESCRIPTION OF PREFERRED EMBODIMENTS
Figures 1 and 2 provide schematic illustrations of the process steps employed
to
process intermetallic alloys from water atomized powders and gas atomized
powders,
respectively. Atomization techniques were used to obtain powders of FeAI alloy
of either
spherical or irregular shape with the target compositions (weight % ) as
indicated in Table 1.
Table 1
Element A1 Mo Zr B C O Fe
Gas atomized 24.0 0.42 0.1 0.005 0.06 0.05 Bal.
Water atomized24.0 0.42 0.1 0.005 0.06 0.31 Bal.
The surface morphologies of gas and water atomized powders are shown in Figure
3.
The water atomization technique yielded irregular shaped powders (Figure 3 A),
and the gas
atomization technique resulted in spherical particles (Figure 3 B).
Unlike the gas atomization technique, water atomization of FeAI alloy powders
required special precautions to reduce the oxygen content of the particles and
prevent
formation of oxides of iron and aluminum on the surface. The oxygen content of
the water
atomized powder is close to 0.3 wt. % or higher. On the other hand, the oxygen
content of the
gas atomized powder is in the range of 0.02-0.04 wt. % , an order of magnitude
lower than the
water atomized powder.
The powders were filled in a steel can and extruded at 1100°C to obtain
fully dense
bars of the alloys. Cylindrical compression specimens of 10 mm diameter and 15
mm height
were machined from the extruded material such that the compression axis is
parallel to the
extruded direction. The flat ends of the specimen had grooves for holding the
lubricant and
the edges were chamfered to avoid initial fold over.
Hot compression tests were conducted in the temperature range 600-
1150°C at 50°C
intervals and constant true strain rate range of 0.001 - 100 s-1 at intervals
of an order of
magnitude. The specimens were allowed to soak at the testing temperature for
about 15
minutes before the start of compression. During testing, the actual
temperature of the
specimen as well as the adiabatic temperature rise were measured using a
thermocouple
inserted in a 1.0 mm hole machined at half its height to reach the center of
the specimen. The
specimens were coated with a borosilicate glass paste that acted not only as a
lubricant but also
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as a protective layer to minimize oxidation. AlI specimens were deformed to
half their height
and air cooled to room temperature following deformation.
The load-stroke data
obtained in compression were processed to obtain true stress-true plastic
strain curves. The
flow stress data obtained at different temperatures, strain rates and strains
were corrected for
adiabatic temperature rise, if any, by linear interpolation between log 6 and
1/T where 6 is the
flow stress and T is the temperature in Kelvin. A cubic spline fit between
log6 and log was
used to obtain the strain
rate sensitivity (m) as a function of strain rate. This was repeated at
different temperatures.
The efficiency of power dissipation (r~) through microstructural changes was
then calculated as
a function of temperature and strain rate using Eq. (2) and plotted as an iso-
efficiency contour
map. The data were also used to evaluate the flow instability parameter ~()
using Eq. (3) as a
function of temperature and strain rate to obtain an instability map.
The starting bar and the deformed specimens were sectioned parallel to the
extrusion
axis and the compression axis respectively and the cut surfaces were prepared
for
metallographic examination.
lyaitial Microstructure
The initial microstructure of the starting rod of water atomized material is
shown in
Fig. 4, which reveals an equiaxed grain structure with an average grain
diameter of about 22
~,m. The large grain size is expected due to the high oxygen content in this
alloy. The prior
particle boundaries are totally eliminated during the extrusion process due to
the possible
occurrence of dynamic recrystallization (DRX). The grains are slightly
elongated in the
extrusion direction which is horizontal in the micrograph. The microstructure
also consists of
a uniform distribution of finer carbide particles and larger alumina
particles.
The microstructure of the starting rod of gas atomized material is shown in
Fig. 5. It
has an equiaxed grain structure with an average grain diameter of about 13
~,m. The prior
particle boundaries are totally eliminated during the hot extrusion process
possibly due to the
occurrence of dynamic recrystallization. The microstructure also consists of
fine carbide and
alumina particles, which are aligned in the direction of extrusion.
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Stress-Strain Behavior
The true stress-true plastic strain curves obtained on samples made from water
atomized powders at temperatures of 600, 650 and 700°C are shown in
Figures 6-~,
respectively. The curves at higher strain rates ( > 1.0 s-1) exhibited peaks
in the flow stress at
strains between 0.1 - 0.2. At the strain rates of 10 and 100 s-1, a second set
of peaks appeared
at larger strains. A visual examination of the specimens showed extensive
cracking early in
deformation. The occurrence of cracks is responsible for the initial drop in
the flow stress,
which then increases when the cracks get welded under compression. At lower
strain rates,
the cracking is less severe due to the compressive state of stress and the
curves indicate some
strain hardening. The cracking problem becomes less severe as the temperature
is increased
and is nearly eliminated at temperatures higher than 700°C.
The true stress-true plastic strain curves obtained on samples made from water
atomized powders in the temperature range 750 - 1150°C are shown in
Figures 9-17,
respectively. These curves may be subdivided into two types depending on the
strain rate. At
strain rates Iower than about 1.0 s-1, the curves are essentially of steady
state type although an
initial drop in the flow stress occurs at temperatures higher than about
950°C. At higher
strain rates, the curves exhibit a peak in the yield stress at a strain, which
is higher at higher
strain rates. For a strain rate of l0 sn the flow reaches a steady state at
larger strains. The
flow stress data obtained from the stress-strain curves at different
temperatures, strain rates
and strains are given in Table 2.
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Table 2. Flow stress data (in MPa) on Iron Aluminide Alloy (water atomized P/M
compact)
obtained in compression at different temperatures, strain rates and strains
(corrected for
adiabatic temperature rise).
StrainStrainTem perature,
C
rate, 750 800 850 900 950 1000 1050 1100 1150
s'
0.001 224.3144.598.8 66.6 53.1 37.1 26.6 12.1 13.0
0.01 330.1245.2168.3114.987.4 65.7 43.3 31.9 23.5
0.1 388.9337.7208.6185.0150.8110.181.8 55.5 44.8
0.1
1.0 475.3423.0355.7276.5245.1190.3131.6102.169.4
10.0 743.2615.3465.1379.3352.4318.2259.4189.4149.4
100.0 840.4751.1555.4455.6383.7354.8313.2264.2174.3
0.001 225.5144.898.5 66.9 53.8 37.3 27.7 13.1 12.3
0.01 331.9246.6169.0115.587.4 66.3 44.5 33.1 22.7
0.1 394.5334.1203.9182.8147.1107.080.5 55.8 43.4
0.2
1.0 492.2436.6355.8272.8238.4184.1127.398.8 68.8
10.0 713.0596.8447.0362.1326.7282.7236.7154.3127.5
100.0 853.0815.4613.4496.5411.8376.1330.9268.8182.5
0.001 226.8144.599.3 68.2 54.3 37.8 28.4 14.1 11.3
0.01 334.8248.9172.3117.088.2 66.4 44.8 33.6 22.0
0.1 395.7330.6205.9183.6147.1107.680.6 56.3 43.0
0.3
1.0 493.6439.9355.3271.1235.6181.4127.498.5 69.4
10.0 677.7603.8461.5373.0341.5291.2230.9151.3121.5
100.0 777.1796.9624.0509.3422.2383.7334.1270.9189.6
0.001 223.3142.498.8 68.2 53.9 38.2 28.8 14.5 10.4
0.01 336.9246.6172.3117.588.1 66.6 44.6 33.7 21.6
0.1 395.4326.0206.6182.0147.7107.180.9 56.5 42.5
0.4
1.0 493.3436.7352.7269.3232.7178.3126.497.2 69.6
10.0 636.8577.3457.8373.5336.9290.0223.5151.7121.3
100.0 683.5753.1605.7501.9421.4382.6327.8261.4189.0
0.001 216.3139.697.4 66.8 54.2 37.7 28.6 14.7 10.2
0.01 335.8242.9170.7115.686.5 65.1 44.0 33.5 21.0
0.1 397.3323.2203.8180.0146.7104.980.5 56.0 41.6
0.5
1.0 497.2431.7347.0266.7228.7174.3123.795.8 68.8
10.0 619.9561.5453.9370.9331.0281:4219.4147.8120.6
100.0 607.7691.6567.1480.2403.4369.4311.6243.5181.5
The true stress - true plastic strain curves obtained on gas atomized material
at
temperatures between 850 and 1150°C and at different strain rates are
shown in Figures 18-
24, respectively. The flow curves at strain rates lower than about 1.0 s-1 are
essentially of
steady state type, although a stress maximum is present at lower strains. At
higher strain
rates, the flow curves exhibit significant flow softening after reaching a
peak in the flow stress
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in the initial stages of deformation. Further, at temperatures higher than
1000°C and at a
strain rate of 1.0 sn, the flow curves exhibit oscillations after the initial
peak in the flow
stress. The flow stress data obtained from the stress-strain curves at
different temperatures,
strain rates and strains are given in Table 3.
Table 3. Flow stress data in MPa for FeAI alloy (Gas Atomized P/M compact)
obtained in
compression at different temperatures, strain rates and strains (corrected for
adiabatic
temperature rise).
Strain StrainTemperature,
C
rate, 850 900 950 1000 1050 1100 1150
s'
0.001 134.0 102.9 62.1 43.1 30.8 22.7 16.2
0.01 188,4 145.2 104.4 73.3 53.5 36.9 27.7
0.1 256.5 215.9 153.1 129.7 96.5 62.6 43.3
0
1
. 1.0 370.6 282.7 255.3 213.8 211.3 109.7 80.9
10.0 551.6 444.4 427.2 338.2 284.5 210.5 152.6
100.0 766,5 566.8 409.7 385.6 330.4 260.4 210.2
0.001 132.6 97.9 62.9 44.9 32.7 23.9 16.5
0.01 188.9 146.1 103.8 73.0 53.5 37.8 28.4
0.1 253.0 213.9 153.8 127.1 94.5 62.4 42.8
0
2
. 1.0 366.9 276.4 247.9 204.9 204.2 106.3 79.1
10.0 590.8 443.4 410.6 316.5 261.0 194.4 153.4
100.0 857.7 609.0 429.6 394.7 336.4 266.4 223.0
0.001 132.8 95.6 63.4 45.8 33.4 24.7 16.4
0.01 189.4 147.7 104.8 71.4 53.7 37.7 29.0
0.1 252.1 214.0 148.4 126.7 93.9 62.4 44.0
0
3
. 1.0 366.8 269.1 243.6 201.8 200.9 105.9 79.0
10.0 592.9 437.5 404.1 306.5 251.2 185.5 153.7
100.0 840.4 605.5 434.3 392.8 333.1 266.5 228.5
0.001 128.0 90.9 62.1 45.5 33.9 24.8 16.0
0.01 185.8 146.1 104.9 68.1 52.6 36.9 28.6
0.1 245.9 213.8 142.2 126.5 92.9 62.1 44.1
4
0
. 1.0 359.0 265.7 237.9 197.5 197.5 103.5 77.7
10.0 561.0 428.6 397.3 295.2 242.4 180.5 147.3
100.0 781.6 585.2 431.5 383.9 324.5 255.7 225.9
0.001 124.1 87.4 61.2 44.7 33.5 24.6 15.7
0.01 181.5 141.7 102.6 64.6 51.9 35.7 28.5
0.1 246.5 209.0 135.9 123.7 91.1 61.7 43.8
0
. 1.0 351.9 259.1 231.8 191.7 193.1 100.7 75.8
10.0 511.9 408.7 384.4 277.6 230.1 169.1 134.9
100.0 725.3 550.1 417.5 364.5 308.5 242.7 217.7
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Comparing and contrasting the true stress-true strain curves for both water
and gas
atomized samples, one sees that the curves are similar. However, the gas-
atomized material
was harder than the water atomized, as may be expected from its finer grain
size. It is not
possible to arrive at the mechanism of hot deformation directly from the
shapes of the stress -
strain curves alone, although they may be correlated with the mechanisms if
evaluated by other
modeling methods. This is because several mechanisms may lead to similar
shapes of stress -
strain curves e.g. flow softening may suggest dynamic recrystallization (DRX),
flow instability
or globalarization of lamellar structures.
WateY Atomized Powders
(i) kinetic analysis
In the hot deformation of materials, the relationship between the steady state
flow
stress, the temperature and strain rate is expressed by a kinetic rate
equation (Eq. 1). In order
to evaluate the stress exponent, the flow stress data obtained at a strain of
0.5 are plotted
against the strain rate on a log-log scale at different temperatures (see Fig.
25). The value of
n, which is the inverse of the slope of the line, is dependent on temperature
being lower at
higher temperatures. Although the data may be fitted to a straight line, it is
clearly seen that
the variation follows a curve, particularly at temperatures lower than
1050°C, indicating that
the value of the stress exponent is strain rate dependent. Ignoring the strain
rate dependence
of the stress exponent, a simple activation analysis yields an apparent
activation energy of
about 465 kJ/mole for the gas atomized material and 430 kJ/mole for water
atomized material.
These values are in agreement with those previously reported (See e.g. J. D.
Whittenberger,
Mater. Sci. Eng. A, 57, 77 (1983) and J. D. Whittenberger, Mater. Sci. Eng. A,
77, 103
(1986)), and are much higher than that for diffusion of either Fe or A1 in
FeAI (See e.g. M.
Eggersmann, et al., Defect and Diffusion Forum, 339, 143 (1997) and E.
Kentzinger, et al.,
Phys. Condens. Matter, 8, 5535 (1996)).
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(ii) processing ~raaps
The processing maps developed at strains of 0.1, 0.2, 0.3, 0.4 and 0.5, on the
basis of
the flow stress data given in Table 2, are shown in Figs. 26-30, respectively.
These maps are
obtained by a superimposition of the instability maps over the power
dissipation maps. In
each of the processing maps, the regime of flow instability as predicted by
the continuum
criterion, given by Eq. (3), is delineated by a thick line (marked as "0")
cutting across several
efficiency contours belonging to the power dissipation maps. The power
dissipation maps
show isoefficiency contours which represent the relative rate of entropy
production occurring
in the material due to microstructural dissipation. They can also be termed as
"microstructural
trajectories" since they actually represent the rate of change of
microstructure occurring during
hot deformation. It is interesting to note that the curvature of the
trajectories changes when
the temperature is increased beyond about 950°C, the temperature at
which dissolution of fine
carbide particles is likely to occur. Since the material system undergoing hot
deformation is
dynamic, non-linear, dissipative and irreversible, it possesses the
characteristics similar to
those exhibiting "deterministic chaos". As the system moves towards a steady
state at higher
strains, the trajectories move towards microstructural attractors (domains of
maximum rate of
entropy production or basins of lowest dissipative energy) depending on their
initial
conditions. These concepts are applied to the materials system for
interpreting the maps
obtained at different strains.
The map obtained at a strain of 0.1 (Fig. 26) exhibits only one domain at a
temperature
of 1100°C and a strain rate of 0.03 s-1 with a maximum efficiency of
power dissipation of
about 44 % . As the strain increases, this domain gives rise to another domain
with a higher
efficiency (i.e. 48% at a strain of 0.2 and increasing with increasing strain,
see Figs. 27-30)
occurring at a temperature of 1150°C and a strain rate of 0.001 s-1.
This domain has a
temperature range of 1000 - 1150°C and a strain rate range of 0.001 -
0.1 s-1.
Simultaneously, another domain also appears beyond a strain of 0.3 with a peak
efficiency of
about 38% occurring at 1150°C and a strain rate of 100 sn. Thus, the
maps given in Figs. 26-
30 suggest that the microstructure of the material is evolving during
deformation as is
commonly observed in systems with several state-space parameters. For example,
the state-
space parameters in the present case are temperature, strain rate, strain and
the rate of entropy
production (dissipative energy state). The best way of representing such a
change is through a
bifurcation diagram.
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Referring to the map obtained at the strain of 0.1 (Fig. 26), the single
domain observed
has a high efficiency of power dissipation which suggests dissipative
mechanisms like DRX.
The efficiency hill representing these contours is not a steep one since an
efficiency range of
only 8 % is spread over a wide temperature range (250°C). Stress-strain
curves in the
temperature range corresponding to this domain exhibit a small drop in yield
stress suggesting
that some significant softening mechanism operates initially. These features
suggest that this
domain represents DRX process and the changes in the maps that occur with
strain give
further support to this interpretation since it causes a change in the
microstructure such that its
response to the temperature and strain rate will change significantly. In high
stacking fault
energy metals, the maximum efficiency of power dissipation for DRX is about 50
% while it is
about 35 % in low stacking fault metals. Thus the observed value of 44 %
suggests that this
iron aluminide alloy has a medium stacking fault energy.
At higher strains (e.g. 0.4, Fig. 29), the processing map exhibits two
domains. The
domain at strain rates below about 0.1 s-1 and at temperatures above
1000°C, has a peak
efficiency of about 56 % and the contours represent a steep hill (efficiency
increases by about
18% within a temperature range of about 125°C). The maximum efficiency
corresponds to a
strain rate sensitivity of about 0.4 and the stress-strain curves are of
steady-state type (Figs.
14-17). Such domains suggest the occurrence of superplastic deformation or
edge cracking of
the material. The results from tensile tests on similar materials have clearly
shown that
abnormal elongations ( > 300 % ) are obtained under these conditions, thereby
confirming the
occurrence of superplasticity in this domain. Typical microstructures obtained
on specimens
deformed at 1100 and 1150°C and strain rates of 0.001 and 0.1 sn are
shown in Fig. 31.
These exhibit very fine equiaxed grain structure. The measured average grain
diameter is
about 12 ~,m, which is much finer than the initial grain size (22 ~,m). Also
the grain size did
not vary significantly with temperature or strain rate as is expected to
happen in the
superplasticity domain.
The higher strain rate domain ( > 10 su) is probably not fully developed
within the
testing regime of temperature and strain rate since only a small part of it is
seen in the map
(Fig. 30). The stress-strain curves under conditions within the domain (Figs.
16 and 17)
exhibit typical DRX features which include a peak in the flow stress followed
by a steady state
as well as initial oscillations reaching a steady state when the strain rate
is at the lower end of
the domain. Typical microstructure recorded on a specimen deformed at
1150°C and 100 s 1
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(Fig. 32) exhibits wavy or irregular grain boundaries, which are considered
signatures of a
DRX process. Thus the high strain rate domain may be interpreted to represent
a DRX
process.
Representative microstructures recorded on specimens deformed at X50°C
at a strain
rate of 0.001 s-1, and at 900°C and 0.1 s-1 are shown in Figs. 33 a-b.
These temperature and
strain rate combinations correspond to regions in which the microstructural
changes are
associated with trajectories that do not get attracted to any of the domains
discussed above.
No significant change is observed in these micrographs. This is further
confirmed by grain
size measurements as plotted in Fig. 34. The grain refinement at temperatures
higher than
1050° may be observed.
A bifurcation diagram at a temperature of 1150°C representing the
changes in the
deformation mechanisms occurring with strain is shown schematically in Fig. 35
which will
help in understanding the changes occurring in the material with strain. Up to
a strain of
about 0.1, dynamic recrystallization occurs in the strain rate range of 0.001
to 1.0 s-1 and
causes grain refinement. At a strain of 0.2, the lower strain rate branch of
the bifurcation
finds superplastic deformation as an attractor when deformed in the strain
rate range about 10-5
s 1 (extrapolated as a mirror reflection) to 10-1 S-1, which continues on
further straining. The
higher strain rate branch of the first bifurcation leads to a DRX attractor
only after a strain of
about 0.2 (critical strain for DRX). This bifurcation occurs in the strain
rate range of 10 to
103 su (the higher strain rate value is an extrapolated one on the basis of a
mirror reflection of
the domain).
The material exhibits flow instabilities at lower temperatures and higher
strain rates as
shown by the instability limit in Fig. 30. These instabilities manifest as
adiabatic shear bands
which are intense at lower temperatures and higher strain rates and flow
localization under
other conditions. Typical microstructures of specimens deformed at
750°C and at three strain
rates in the instability region are shown in Figs. 36 a-c. An intense
adiabatic shear band with
associated cracking is recorded in the first case and flow localization is
more diffused in
others. These conditions may be avoided in processing this material.
(iii) desigfa of hot woYki~g pYOCess
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On the basis of the constitutive behavior of iron aluminide alloy as revealed
in the
processing maps, the following hot working schedules may be designed for the
bulk working
of the material.
Since the material exhibits a change in the mechanism of hot deformation at
strains
above 0.1, it is beneficial to "condition" the material by hot working it at
1100°C and at strain
rates in the range 0.001 - 1.0 s-i using small strains. This may be done
either by forging or
rolling. Once the billet is conditioned, the material has extensive
workability at temperatures
above 1100°C both at higher strain rates (10 s-1) due to DRX and lower
strain rates ( < 0.1 s-1)
due to superplastic deformation. The higher strain rate domain may be
exploited for
continuous rolling of the material since this process is generally done at
higher speeds.
However, component manufacture from the sheets may be done by superplastic
forming
associated with diffusion bonding with a suitable material. Manufacture of
other forged
components are best done in the high strain rate domain using processes
including drop
forging which is a cost effective process.
Gas Atomized Powders
(i) kinetic analysis
Within the domain of DRX (in the
temperature range 950 - 1150°C and strain rate range 0.001 - 1.0 s-1),
kinetic analysis using
Eq. (1) has been conducted. Considering the variation of log (6) vs. log () in
the above
limited temperature and strain rate ranges to be approximately linear, the
value of the stress
exponent is estimated to be about 4.4 (an average strain rate sensitivity of
about 0.23). The
Arrhenius plot giving the variation of log (6) with (1/T) is shown in Fig. 37,
from which an
apparent activation energy of about 465 kJ/mole has been estimated for the
process of DRX in
this material. This value is higher than that for the diffusional processes in
FeAI, as is
commonly observed for DRX.
It is customary to correlate the grain size variations in the DRX domain with
the
Zener-Hollomon parameter, Z, given by:
2=E exp [Q~RT] (4)
Such a variation is shown in Fig. 38, which is linear as expected for the DRX
process. Such
plots are useful in controlling grain size in the material during processing.
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(ii) processirzg maps
The processing maps developed at strains 0.1, 0.2, 0.3, 0.4 and 0.5, on the
basis of the
flow stress data given in Table 3, are shown in Figs. 39-43, respectively.
These maps are
obtained by a superimposition of the instability maps over the power
dissipation maps. In
each of the processing maps, the regime of flow instability as predicted by
the continuum
criterion given by Eq. (3) is delineated by a thick line (marked as "0")
running across several
efficiency contours belonging to the power dissipation maps. The power
dissipation maps
show isoefficiency contours, which represent the relative rate of entropy
production occurring
in the material due to microstructural dissipation. They can also be termed as
"microstructural
trajectories" since they actually represent the rate of change of
microstructure occurring during
hot deformation. The curvature of the trajectories changes when the
temperature is increased
beyond about 875°C, the temperature at which dissolution bf fine
carbide particles is likely to
occur. Since the material undergoing hot deformation is a non-linear
dissipator of power, the
microstructural trajectories get attracted to basins of lower dissipative
energy and form
domains where the efficiency of power dissipation is maximized. These domains
represent
specific microstructural mechanisms, which may be identified by metallographic
examination
of specimens deformed in this domain.
Referring to the map obtained at the strain of 0.1 (Fig. 39), the single
domain observed
has a peak efficiency of power dissipation of 44% occurring at about
1075°C and 0.1 s n.
Maps obtained at higher strains (Figs. 40-43) are not significantly different
from that obtained
at a strain of 0.1, although the peak efficiency of the domain referred to
above has slightly
decreased from 44 % to 40 % . In high stacking fault energy metals, the
maximum efficiency of
power dissipation for DRX is about 50 % while it is about 35 % in low stacking
fault energy
metals. Thus the observed value of 40 - 44% suggests that this iron aluminide
alloy has a
medium stacking fault energy.
Typical microstructures obtained on specimens deformed at a strain rate of 0.1
s n and
at temperatures of 1050 and 1100°C are shown in Figs. 44 a-b. The
microstructures exhibit
fine equiaxed grain structure with irregular grain boundaries typical of
dynamic
recrystallization. The variation of average grain diameter with deformation
temperature is
shown in Fig. 45, which exhibits a sigmoidal curve typically observed when
dynamic
recrystallization occurs. The variation of the efficiency of power dissipation
with temperature
as obtained from the processing map (Fig. 43) at the strain rate of 0.1 s-1 is
also shown in
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Fig. 45. It may be noted that the temperature at which 50 % variation in grain
size has
occurred coincides with the temperature for the peak efficiency in the domain
(called DRX
temperature). This is also a typical feature of the DRX process. At
deformation temperatures
lower than the DRX temperature, the grain size gets finer (10 ~,m or less)
than the initial value
(13 ~cm) while it increases (to about 20 ,um) at higher temperatures. The
microstructures of
the material deformed at 1100 ° C and at strain rates of 0.01 and 1.0 s
a are shown in Figs . 46
a-b and the variation of grain size with strain rate in the domain is plotted
in Fig. 47, both of
which show that the grain size is finer at higher strain rates and the
variation is linear with
respect to log strain rate. At the optimum temperature and strain rate for
DRX, the
workability of the material is maximum. Thus the hot working processes may be
designed for
workability optimization as well as for microstructural control within this
domain of DRX.
As per the instability criterion given by Eq. (3), the material exhibits flow
instabilities
at strain rates higher than 10 s-1 in the temperature range 950-1100°C.
These instabilities
manifest as bands of flow localization as seen in the microstructure of the
specimen deformed
at I00 s 1 and 1050°C given in Fig. 4~. The conditions of instability
predicted by Eq. (3) may
be avoided in processing this material.
(iii) design of hot working pYOCesses
On the basis of the constitutive behavior of iron aluminide alloy as revealed
in the
processing maps, the following hot working schedules may be designed for the
bulk working
of the material.
Since the material undergoes dynamic recrystallization in the temperature
range 950 -
1150 °C and strain rate range 0.001 - 1.0 sn, with a peak efficiency of
power dissipation
(44%) occurring at 1075 °C and 0.1 s-1, these parameters represent the
optimum conditions
for working this material. Manufacture of components is best done in this low
to moderate
strain rate domain using processes such as extrusion and press forging.
Conversely, at high strain and high strain rates at temperatures between 950
and 1000
°C, flow instabilities occur, which should be avoided. This suggests
that this alloy is not
optimized for processing by, for example, continuous rolling, which is
generally done at
higher speeds, or drop forging, which involves high strain rates.
Additionally, the grain size varies sigmoidally with temperature at the strain
rate
corresponding to peak efficiency in the DRX domain while the grain size
decreases with strain
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rate in a linear fashion. Therefore, control of both the temperature and
strain rate
environment is important to obtain the desired microstructure. Further, the
selection of this
environment may be aided by the use of the Zener-Hollomon parameter.
The difference in the behavior of the gas atomized and water atomized powder
compacts is attributed to the reduced oxide particle content in the former
case.
Structural ayad ElectYical Resistance Applications
The FeAI based alloys and the processing methods developed here are intended
for use
in industrial and domestic applications. Some examples of possible uses
include as heat
treatment and furnace fixtures in the thermal processing industry, as heating
elements and
resistance alloys, and as forged components such as automotive valves. In
these applications,
the superior corrosion resistance of FeAI based alloys coupled with the
reduction in
manufacturing costs are attractive.
Although the present invention has been described in connection with preferred
embodiments thereof, it will be appreciated by those skilled in the art that
additions, deletions,
modifications, and substitutions not specifically described may be made
without department
from the spirit and scope of the invention as defined in the appended claims.
