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In-Situ Reactive Processing of Nickel Aluminides
by Laser-Engineered Net Shaping
WEIPING LIU and J.N. DUPONT
Nickel aluminide intermetallics (e.g., Ni3Al and NiAl) are considered to be attractive materials for
high-temperature structural applications. Laser-engineered net shaping (LENS) is a rapid prototyping
process, which involves laser processing fine metal powders into three-dimensional shapes directly from
a computer-aided design (CAD) model. In this work, an attempt has been made to fabricate aluminide
intermetallic compounds via reactive in-situ alloying from elemental powders using the LENS process.
In-situ reactive alloying was achieved by delivering elemental Ni and Al powders from two different
powder feeders, eliminating segregation observed in the samples deposited by using the premixed
elemental powders. Nickel aluminides of various compositions were obtained easily by regulating the
ratio of their feed rates. The aluminide deposits exhibited a high solidification and subsolidus cracking
susceptibility and porosity formation. The observed porosity resulted from a water-atomized Ni powder
and can be minimized or eliminated by the use of a N2-gas-atomized Ni powder of improved quality.
Cracking was due to the combined effect of the high thermal stresses generated from the LENS processing
and the brittleness of the intermetallics. Crack-free deposits were fabricated by preheating the substrate
to a temperature of 450 °C to 500 °C during LENS processing. Compositionally graded Ni-Al deposits
with a gradient microstructure were also produced by the in-situ reactive processing.
I. INTRODUCTION
NICKEL aluminide intermetallics (e.g., Ni3Al and NiAl)
are considered to be attractive materials for high-temperature
structural applications, for example, as turbine elements or
other heat- and oxidation- or wear-resistant components.[1,2,3]
This is due to their high strength retention at elevated
temperatures, combined with relatively low density and good
oxidation and corrosion resistance. For example, the elevatedtemperature strength and creep resistance of recently developed
Ni3Al alloys are shown to be superior to most commercial
superalloys.[2] The high-temperature corrosion and oxidation
resistance of these aluminide intermetallics also makes them
good candidates for applications as high-temperature coatings.
In spite of many excellent physical and chemical properties,
these materials suffer from low ductility and toughness at
ambient temperature and, therefore, are difficult to fabricate
by conventional processing methods. Near-net-shape fabrication of the intermetallics is possible by casting and powder
metallurgical processes. However, the processing costs are
high, and product shapes are usually limited because of the
use of molds or dies in these techniques.
Laser-engineered net shaping (LENS) is a solid freeform
fabrication process, which involves laser processing fine metal
powders into fully dense three-dimensional shapes directly
from a computer-aided design (CAD) model. The LENS
process is able to fabricate complex prototypes in near-net
shape, leading to significant time and machining cost savings.
It also has potential for precision repair, fast tooling, and smalllot production.[4] So far, this technique has been used in the
laboratory. A variety of metals and alloys have been deposited
by LENS processing, such as H13 steel, 316 stainless steel,
WEIPING LIU, Visiting Scientist, and J.N. DUPONT, Associate
Professor, are with the Department of Materials Science and Engineering,
Lehigh University, Bethlehem, PA 18015. Contact e-mail: [email protected]
Manuscript submitted December 9, 2002.
METALLURGICAL AND MATERIALS TRANSACTIONS A
nickel-based superalloys, and titanium alloys.[5–8] However,
most of these deposits were fabricated using a prealloyed
powder as the feedstock material, although Schwendner et al.[9]
investigated the deposition of Ti-10Cr and Ti-10Nb alloys from
a blend of elemental powders using the LENS process. In
addition, little work has been done on processing aluminide
intermetallics by LENS, although a few articles have been
published on microstructural characterization of laser-deposited
TiAl alloys fabricated from prealloyed titanium aluminide
powders.[10,11]
In this study, an attempt has been made to fabricate
aluminide intermetallics via reactive in-situ alloying from a
blend of elemental powders using the LENS process. The
in-situ reactive fabrication of intermetallics via LENS has
the following potential advantages: (1) raw-material cost
savings by eliminating the fabrication steps required for
prealloyed powders; (2) suitability for fabricating a compositionally graded structure and materials (e.g., a Ni-Ni3AlNiAl compositionally graded material); and (3) energy savings
by use of the reaction-generated heat.
II. EXPERIMENTAL PROCEDURE
The powders used in this investigation were air-gas-atomized Al (Grade 120, purity 99.7 pct) from Alcoa (Rockdale,
TX), and water-atomized Ni (Ni-118, purity 99.6 pct) from
Praxair Surface Technologies (Indianapolis, IN). A nitrogengas-atomized Ni powder from Crucible Research (Pittsburgh,
PA) was also used in later experiments to resolve a porosity
problem associated with the water-atomized Ni powder. No
boron doping was used in this study, although it is known
that boron can suppress the brittleness of unalloyed Ni3Al.
The Ni-118 and Grade 120 Al powders had a mesh size of
200/325 (particle sizes between 45 and 75 m), while
the N2-gas-atomized Ni had a mesh size of 100/325
(particle sizes between 45 and 150 m). A 6.35-mm-thick
pure nickel plate (Ni 200) was used as the substrate material.
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The substrate was ground with 320-grit SiC paper and then
grit blasted before deposition.
An Optomec LENS 750 system was used in this study to
deposit the samples, which have a size of 12.7 12.7 mm 6 to 8 layers. The LENS machine consists of a Nd:YAG laser,
a four-nozzle coaxial powder feed system, a controlledenvironment glove box, and a motion control system. The Nd:
YAG laser has a 0.5- to 1-mm-diameter circular beam at the
focal zone, with the Gaussian intensity distribution and a
maximum output power of 750 W. The powder-delivering
nozzles are designed and arranged in such a way that the
powder streams converge at the focal point of the laser beam.
To make a LENS deposit, a CAD model of a three-dimensional
component is first sliced into a series of layers of finite thickness
using computer software. Each of these layers is then translated
into a series of line patterns in order to deposit the layers.
The laser beam is used as a heat source to create a molten pool
on the substrate, and the powder is injected into the melt pool
by an inert gas flowing through the powder feed system. The
first layer of the component is bonded to the substrate. The
substrate, together with the component under fabrication, is
moved along the line patterns in the plane of current layer with
the motion control system. After completing a layer, the laser
focal point and powder-delivering nozzles are incremented
upward in the height direction in an amount of the layer
thickness. A new layer is subsequently deposited onto the
previous layer until the component is fully constructed in
the layer-by-layer fashion. The substrate can be removed when
the final product is fabricated. The LENS processing is performed in an argon atmosphere in the controlled-environment
glove box to prevent oxidation.
In the initial experiments of this work, Ni and Al powders
were mixed using jar milling before being loaded into the
single powder feeder. It was found that homogeneous mixing
could not be achieved due to a large difference in powder
density. Therefore, two powder feeders were used in all the later
experiments to deliver the Ni and Al powders separately. In
this case, the Ni and Al elemental powders were mixed
in situ during feeding from the powder feeders to the nozzles,
and their feed rates were controlled individually by regulating
the rotational speed of the powder feeder. In the present
investigation, successive layers are deposited with the bead
lines of two adjacent layers at an angle of 90 deg. The laser
power and traverse speed were chosen according to the Ni and
Al powder feed rates in the experiments. In order to realize
the advantages of the in-situ reactive processing for functionally
graded materials, compositionally graded Ni-Al deposits were
fabricated in this study. For this purpose, the rotational speed
of the powder feeder for Al varied from zero to a preset value
stepwise as the number of layers increased, while the rotational
speed of the powder feeder for Ni was reduced stepwise from
a maximum value which was set according to the processing
parameters used to another preset value. The variations of rotational speeds for the powder feeding were program controlled,
with a minimal step of 0.1 rotations per minute (rpm) used in
this work. In all the LENS experiments, the layer thickness,
hatch spacing, and stand-off distance were set at 0.254, 0.381,
and 152.4 mm, respectively. The oxygen level in the glove
box was kept below 10 ppm during processing.
Wet chemical analyses were conducted for selected
deposit samples. Light optical microscopy (LOM) and X-ray
diffraction (XRD) were used for microstructure and phase
2634—VOLUME 34A, NOVEMBER 2003
analysis in the study. Samples for LOM were mounted and
polished using standard metallographic techniques and
etched with the Marble’s reagent (10 g CuSO4, 50 mL HCl,
and 50 mL H2O). The XRD was conducted on the sample
surface perpendicular to the build direction. The surfaces
of cracks were examined with a JEOL* 6300 scanning
*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.
electron microscope (SEM). Microhardness measurements
were conducted on the compositionally graded Ni-Al
deposits using a Vickers indenter under a 300-g load.
III. RESULTS AND DISCUSSION
A. Role of Exothermic Reactions
In the LENS process, the powders are fed into the laser
focal zone and melted in the molten pool. Therefore, the
Ni and Al powder particles are in a liquid state before the
synthesis reaction takes place. However, it should be pointed
out that there exists the situation where the Ni and Al
particles are not both completely melted before forming a
nickel aluminide. This happens only when the processing
parameters are not properly chosen (e.g., very high powder
feed rate and low laser power or high travel-speed combinations). It is usually accompanied by a varying degree of
porosity in the product (as in the case of conventional
combustion synthesis processes), and can be avoided easily
by adjusting the processing parameters.
There are five intermetallic compounds (Al3Ni, Al3Ni2,
Al3Ni5, NiAl, and Ni3Al) existing in the Ni-Al binary-phase
diagram (Figure 1).[12] Each of these intermetallics has a
negative heat of formation. Of these intermetallics, Ni3Al and
NiAl are technologically the most significant and have
attracted much scientific interest. Their enthalpies of formation
at 298 K are 153.1 and 118.4 kJ/mol, respectively.[13]
Generally, for the formation of a melt of nickel aluminide
NiaAlb (l) via the reaction
aNi(l) bAl(l) → NiaAlb(l) heat
[1]
where “l” in the parentheses denotes a liquid state. The
following thermal-balance equation is applicable on the
adiabatic condition:
aHNi bHAl HNiaAl b Hf,298(NiaAlb )
[2]
where
HNi Tm,Ni
Cp, s (Ni)dT Hm,Ni 298
Tm, Al
HAl Tr,Ni
Cp, l (Ni)dT [3]
Tm,Ni
Tr, Al
Cp, s (Al)dT Hm,Al 298
HNia Al b Cp, l(Al)dT
[4]
Tm, Al
Tm,Nia Alb
Cp, s (NiaAlb ) dT Hm,Nia Al b
[5]
298
Tad
Cp,l (NiaAlb )dT
Tm, NiaAlb
and
Hf,298(NiaAlb) the enthalpy of formation of NiaAlb at
298 K,
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Cp,s (substance) the specific heat capacity of the
substance in the solid state,
Cp,l (substance) the specific heat capacity of the
substance in the liquid state,
Hm,Ni the enthalpy of fusion for Ni,
Hm,Al the enthalpy of fusion for Al,
Hm,Nia Alb the enthalpy of fusion for NiaAlb,
Tm,Ni the melting temperature of Ni,
Tm,Al the melting temperature of Al,
Tm,NiaAlb the melting temperature of NiaAlb,
Tr,Ni the starting temperature of reaction for Ni liquid,
Tr,Al the starting temperature of reaction for Al liquid,
and
Tad the adiabatic temperature of the reaction (the
maximum attainable temperature).
From Eqs.[1] through [5], the adiabatic temperature can
be calculated. Using the thermodynamic data[13,14] provided
in Table I and assuming the starting temperatures of reaction
for Ni and Al liquids to be 1728 and 933 K, respectively,
the adiabatic temperature of the reaction was obtained to be
2815 K for Ni3Al (1) and 2703 K for NiAl (1). Compared
with the starting temperatures of the reactant liquids, this
represents a significant increase in temperature due to the
exothermic reaction of the liquid Al and Ni. Deevi and
Sikka[15] observed that the addition of an aluminum melt
stock into a molten nickel for obtaining a Ni 3Al melt
resulted in a sudden increase in temperature from 1600 °C
(1873 K) to 2300 °C (2573 K) and above. Their observation
provided a reasonable support for the previously calculated
adiabatic temperature of the reaction. It should be pointed
out that this calculated temperature increase is only limited
to the reaction zone (not the whole melt pool) under the
adiabatic condition for the LENS process. The actual
temperature increase in the melt pool due to the exothermic
reaction should be much lower, because of the larger
volume of previously formed nickel aluminide melt in the
pool and because of the heat loss in the process as well.
As mentioned earlier, in the LENS process, a focused
laser beam is used as the heat source to melt the powders.
The reaction-generated heat is, therefore, not critical for
the melting process, although it can reduce the required laser
energy. The exothermic reactions, however, are very important for the in-situ synthesis (alloying) process. According
to the results of research on self-propagating high-temperature
synthesis (SHS), the reaction velocity (the velocity of the
propagating combustion wave) in SHS can be extremely
high. In the case of the Ni Al reaction, the melting of Ni
plays an important role in the temperature dependence of
the reaction velocity, which can reach more than 10 cm/s
in the liquid state.[16] Using this velocity and the maximum
particle size of 150 m used in the present study, the time
taken for full reaction (fully intermixing) can be estimated
to be 1.5 103 s. Therefore, the liquid nickel aluminide
(NiaAlb) can be instantly synthesized in the melt pool once
the Ni and Al liquids meet during LENS processing. The
reaction velocity is much higher than the travel speeds (Vb)
used in the present study and, consequently, is also much
higher than the solidification velocities (Vs) of the molten
metal in the melt pool, according to Vs Vb . cos , where
is the angle between the solidification-front normal
and the travel direction of the melt pool. This ensures the
homogeneity of in-situ alloying and the resulting intermetallic
compounds.
B. In-Situ Reactive Alloying and Characterization of the
Deposits
Nickel aluminides of various compositions were deposited
by changing the feed rates of the two powder feeders
individually, which were implemented by regulating the
rotational speed (RPM) of each powder feeder. Based on the
chemical-analysis results, it was found that the composition
of in-situ synthesized nickel aluminide deposits depends essentially on the ratio of the RPM of the two powder feeders,
Fig. 1—Ni-Al binary phase diagram.[12]
Table I. Thermodynamic Data[13,14] Used for the Calculation of Adiabatic Temperatures of the Exothermic Reactions
Substance
Tm (K)
Hf,298 (kJ mol1)
Hm (kJ mol1)
Specific Heat Capacity (J mol1K1)
Cp,s 32.60 1.80 103T 5.58 105T 2 (T 298 to 630 K)
Cp,s 29.68 4.18 103T 9.32 105T 2 (T 630 to 1728 K)
Cp,l 38.87
Cp,s 20.65 12.37 103T
Cp,l 31.77
Cp,s 88.49 32.22 103T 0.001 105T 2 0.001 106T 2
Cp,l 142
Cp,s 41.93 13.6 103T 0.033 105T 2 0.1 106T 2
Cp,l 71
Ni
1728
0
17.6
Al
933
0
10.9
Ni3Al
1668
153.1
50
NiAl
1911
118.4
63
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regardless of their individual values. Figure 2 shows the chemical-analysis results of Ni and Al in the deposited samples
vs the ratio of the rpm for Al to the rpm for Ni used in the
LENS deposition. As can be seen, a regressed linear relationship exists between the two under the present experimental
conditions. This plot is useful for choosing the appropriate
RPM of the powder feeders to obtain a deposit of desired
compositions.
Figure 3 shows the XRD patterns of the selected nickel
aluminide deposits of various compositions. These results
indicate that both single-phase and two-phase nickel aluminide alloys have been in-situ alloyed from the elemental
(b)
(a)
Fig. 2—(a) and (b) Relationship between the chemical compositions of the in-situ alloyed Ni-Al deposits and the ratio of rotational speeds (rpm) of the
two powder feeders for Al and Ni, respectively.
Fig. 3—(a) through (d) X-ray diffraction patterns of nickel aluminide deposits of various compositions.
2636—VOLUME 34A, NOVEMBER 2003
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nickel and aluminum powders via LENS processing. The
deposit of a composition 76.4Ni-23.5Al (in wt pct, hereafter)
is a single-phase alloy consisting of Ni1.1Al0.9, while the
87.9Ni-10.4Al, 80.4Ni-18Al, and 64.6Ni-35.3Al deposits
are all two-phase alloys. The 64.6Ni-35.3Al deposit consists
of Ni0.9Al1.1 and Ni2Al3, and the 80.4Ni-18Al deposit consists
of Ni3Al and Ni1.1Al0.9 phases. Ni3Al and possibly some
-Ni (nickel solid solution) are present in the 87.9Ni-10.4Al
deposit. According to the Ni-Al binary-phase diagram
(Figure 1), the XRD results agree relatively well with what
is predicted from the equilibrium phase diagram, except that
Ni2Al3 phase is present in the 64.6Ni-35.3Al deposit due
to rapid solidification conditions associated with the LENS
processing.
Figure 4 shows the LOM micrographs of a Ni3Al deposit.
Irrespective of the gas porosity and cracks, which will be
discussed in the following sections, the observed macrostructure is essentially columnar, with the axis being approximately parallel to the build direction (Z direction), indicating
a very high temperature gradient existing in this direction.
Generally, relatively small grains with a low aspect ratio are
observed near the deposit-substrate interface due to epitaxial
grain growth from the substrate, while larger grains with a
higher aspect ratio are present in subsequently deposited
layers as a result of competitive grain growth. A similar
grain morphology was observed in the LENS-deposited
Fig. 4—LOM micrographs of a Ni3Al deposit: (a) layers close to the
substrate and (b) layers close to the deposit surface.
METALLURGICAL AND MATERIALS TRANSACTIONS A
Ti-6Al-4V samples by Kobryn et al.[7] Because of the nature
of competitive growth, the developed columnar grains are
preferentially oriented with their easy growth direction
100
(for fcc and bcc metals and alloys) being parallel
to the direction of the maximum temperature gradient, i.e., the
build direction of the deposit in the LENS processing. This
explains why the diffractive peak from {200} planes in the
XRD pattern for the Ni3Al deposit (as shown in Figure 3(a))
appears as the strongest one in place of the peak from {111}
planes. Figure 5 shows the microstructures of a Ni3Al and
a NiAl-Ni3Al deposit, respectively. The observed subgrain
structure is cellular/dendrite. The 80.4Ni-18Al deposit has
a peritectic structure consisting of the -NiAl and ’-Ni3Al
phases.
C. Processing Defects, Causes and Solutions
1. Porosity
There are two types of porosity observed in the in-situ
alloyed aluminide deposits: gas porosity (with spherical
shape, as shown in Figure 4) and lack-of-fusion porosity
(with irregular shape, as shown in Figure 6). The lack-offusion pores were always observed along layer boundaries
and are believed to be formed due to insufficient melting
at the interlayer boundaries. This kind of porosity increased
as the Al content was increased in the deposit. There are
Fig. 5—Microstructures of the in-situ alloyed nickel aluminide deposits:
(a) the 87.9Ni-10.4Al deposit (440 W, 5.9 mm/s); and (b) the 80.4Ni-18Al
deposit at the top layer (440 W, 7.6 mm/s).
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Fig. 6—Porosity and cracks observed in the 76.4Ni-23.5Al deposit (black
arrows indicate the lack-of-fusion porosity).
two reasons for this experimental observation. One is that
the Al powder particles have a lower laser absorption
coefficient and, as a result, less energy was absorbed from
the laser beam at a higher Al content in the powder mixture
when the same processing parameters were used. Another
reason is that NiAl was synthesized at a higher Al content,
which has a higher melting point (1638 °C) and needs more
energy to melt than Ni3Al, with a melting temperature of
1395 °C. The lack-of-fusion porosity could be eliminated
by increasing the laser power and/or decreasing the traverse
speed. However, the spherical gas porosity was found to
be a more severe problem. This kind of porosity appeared
more in the layers close to the substrate and less in the
subsequent layers (Figure 4), probably due to a preheating
effect from the previously deposited layers. The preheating
effect leads to a reduced cooling rate and, hence, a longer
solidification time, so that less gas can be trapped in the
solidified microstructure. The present experiments demonstrated that the amount of gas porosity could be decreased
to some extent by decreasing the traverse speed or by baking
the powders at 200 °C for 6 hours under flowing Ar gas
before deposition.
It was also found in the experiments that gas porosity
increased with increasing Ni content in the deposit. Subsequently, the starting Ni and Al powders were mounted and
polished for LOM examination. Figures 7(a) and (b) show
the photomicrographs of as-polished Ni and Al powders,
respectively. As can be seen, the water-atomized Ni powder
particles have a very rough, nodular surface and contain
significant amounts of porosity. The Al powder particles
have a relatively better quality. Therefore, the gas porosity
observed in the Ni-Al deposits can be concluded to result
from two main factors: (1) moisture present at the surfaces
of starting powders, and (2) gas porosity within the Ni
powder particles. Baking the powder before deposition could
remove the surface moisture, but not the gas porosity within
the powder particles. Consequently, the use of gas-atomized
Ni powder with minimal porosity should be the ultimate
solution to the gas porosity problem.
Figure 8(a) shows a micrograph of the nickel aluminide
deposit processed from a nitrogen-gas-atomized Ni powder
2638—VOLUME 34A, NOVEMBER 2003
Fig. 7—Photomicrographs of as-polished cross sections of the starting
powders: (a) Praxair Ni-118 powder and (b) Alcoa grade 120 Al powder.
and the same Al powder. The as-polished cross sections of
the Ni powder particles are shown in Figure 8(b), which
indicates Ni particles of relatively spherical shape with a
smooth and clean surface and only occasional discrete inside
pores. As can be seen in Figure 8(a), the resultant deposit
is much denser (essentially pore-free), due to the use of the
gas-atomized Ni powder of improved quality. This effect
has also been observed in the work conducted on LENS
deposits of stainless steels by Susan et al.[17]
2. Cracks
Another problem encountered in fabricating the in-situ
alloyed nickel aluminide deposits is the occurrence of cracking
in the deposits. Both macro- and microcracks were observed
in the deposits. As can be seen in Figures 4 and 6, cracking
occurred predominantly intergranularly in the nickel aluminide
deposits. It was found that cracking susceptibility increased
at higher Al contents. The cracking susceptibility could be
reduced to some extent by increasing the incident laser heat
input (incident laser power over traverse speed). Figure 9
shows the representative SEM fractographs of the crack
surfaces in an in-situ alloyed NiAl deposit. Fracture occurred
mainly along the solidifying cells/dendrites (grain-boundary
residual liquids). But, solid-state intergranular failure features
were also present, as indicated in Figure 9(a). The fracture
features indicate that the cracking is essentially a type of
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Fig. 8—(a) Micrograph of a Ni-Al deposit fabricated from gas-atomized
Ni and Al powders (410 W, 6.4 mm/s); and (b) as-polished cross sections
of the N2 gas-atomized Ni powder.
hot cracking, i.e., solidification cracking plus subsolidus
cracking. Previous weldability studies[18,19,20] of Ni3Al-based
nickel aluminides also demonstrated a hot-cracking susceptibility in both the weld fusion zone and heat-affected zone.
As is known from welding metallurgy principles, hot
cracking occurs within a brittleness temperature range during
the weld solidification and cooling process. Generally, a wide
solidification temperature range increases the cracking
susceptibility.[21] From the Ni-Al binary-phase diagram, the
Ni-Al alloy with a primary -NiAl phase has a larger
solidification-temperature range, and, thus, a higher propensity for cracking was observed. Another important factor
that leads to cracking in deposits is the thermal stresses
generated during the LENS processing. Hot cracking
occurred as the tensile stress induced by solidification shrinkage and thermal contraction of the deposit exceeded the
material’s fracture strength in the brittleness temperature
range. Since a high level of thermal stresses builds up during
LENS processing and the nickel aluminides are generally
more brittle than conventional alloys, they exhibited a higher
cracking susceptibility. In addition, extensive grain-boundary
microcracking was observed in the Ni3Al deposit (Figure 4).
This might be due to an additional factor: the grain-boundary
brittleness of unalloyed Ni3Al. Boron doping of the alloy
should mitigate this kind of microcracking.
METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 9—(a) and (b) SEM fractographs at different magnifications of hot
cracks observed in an in-situ alloyed NiAl deposit (arrows indicate solidstate fracture sites).
For the purpose of decreasing the thermal stresses and,
thus, preventing the deposit cracking, a hot plate was used
to preheat the substrate to a temperature of 450 °C to 500 °C
during the in-situ reactive LENS processing. Crack-free
nickel aluminide deposits were obtained by processing with
substrate heating, a 410 W laser power, and 6.4 mm/s
traverse speed. Figure 10 shows the photographs of the
as-polished surfaces of the deposits fabricated without and
with substrate heating during processing, respectively. These
deposits were processed from the water-atomized Ni powder
and gas-atomized Al powder. It can be seen that the LENS
processing with substrate heating can produce crack-free
deposits as well as reduce the porosity formation in the
deposits.
D. Compositionally Graded Ni-Al Deposits
On the basis of the previous investigation into the in-situ
reactive processing, compositionally graded Ni-Al deposits
were fabricated by controlling the rotational speeds of the
two powder feeders and using the substrate preheating
technique for preventing cracking. Figure 11 shows the
designed (calculated from the ratio of the rotational speeds)
variation of Al content in a compositionally graded Ni-Al
deposit and the Vickers microhardness distribution along the
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increased from zero to 34 wt pct, the microstructure changed
gradually in the following order: Ni S Ni/Ni3Al S Ni3Al
S Ni3Al/NiAl S NiAl. The microhardness increased
accordingly from approximately 110 Vickers hardness number
(VHN) for Ni to about 410 VHN for NiAl. The NiAl
aluminide has the highest melting point in the Ni-Al system
and exhibits excellent oxidation and corrosion resistance at
high temperatures.[22] The Ni S Ni3Al S NiAl graded
structure is expected to offer potential for high-temperature
applications, for example, as claddings or surface materials
on nickel alloys and steels.
IV. CONCLUSIONS
Fig. 10—As-polished surfaces of the deposits fabricated (a) without and
(b) with substrate heating during LENS processing (rounded white spots
in the photographs are porosity).
A study has been conducted to fabricate nickel aluminide
intermetallics via reactive in-situ alloying from a blend of
elemental powders using the LENS process. In-situ reactive
alloying was successfully achieved by delivering elemental
Ni and Al powders from two different powder feeders. The
rapid exothermic reactions between the liquid Ni and Al
ensure the homogeneity of in-situ alloying and the resultant
intermetallic compounds. The composition of in-situ alloyed
Ni-Al deposits depends, essentially, on the ratio of rotational
speeds of the two powder feeders. Nickel aluminides of
various compositions were fabricated by regulating the ratio
of their feed rates. The nickel aluminide deposits exhibited
a high solidification and subsolidus cracking susceptibility
and porosity formation. The observed high porosity contents
resulted from a water-atomized Ni powder and can be
minimized or eliminated by the use of a gas-atomized Ni
powder of improved quality. Cracking was due to the
combined effect of the high thermal stresses generated from
the LENS processing and the brittleness of the intermetallics.
Crack-free deposits were obtained by preheating the substrate
to a temperature of 450 °C to 500 °C during LENS processing.
Compositionally graded Ni-Al deposits were fabricated by
the in-situ reactive processing, with a gradient microstructure
in the following order: Ni S Ni/Ni3Al S Ni3Al S
Ni3Al/NiAl S NiAl.
ACKNOWLEDGMENTS
The authors gratefully acknowledge support of this work
by the National Science Foundation through a PECASE
Award, Grant No. DMI 9983968, made through the Division
of Manufacturing and Industrial Innovation of NSF.
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Fig. 11—Microhardness distribution and the designed variation of Al content
in a compositionally graded Ni-Al deposit.
thickness direction. The graded deposit was processed with
a 410 W laser power and a 8.5 mm/s traverse speed. The
graded, microstructurally constituent phases are also indicated
at the top of the figure. As the Al content in the deposit
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