1. INTRODUCTION
High-temperature metallic materials are generally developed with consideration for
their performance under severe conditions such as creep, fatigue, hot corrosion, and
thermal degradation. Ni-based superalloys, which are representative high-temperature
metallic materials, exhibit excellent oxidation resistance and mechanical properties
at elevated temperatures; for this reason, they are widely employed as critical components
exposed to harsh environments in aerospace, space, and power-generation applications[1-4]. However, if the alloy composition is not properly controlled, the formation of brittle
phases can occur, leading to a degradation of mechanical properties and oxidation
resistance. Accordingly, extensive research has been conducted to improve the high-temperature
oxidation resistance and mechanical performance of Ni-based superalloys by controlling
the contents of alloying elements such as Cr, Al, and Fe, thereby modifying oxide
scale types, increasing the volume fraction of strengthening phases, and reducing
diffusion rates at elevated temperatures[1-5].
Metallic foams are being utilized in various engineering applications, such as sandwich
structure cores, catalyst substrates, filters, and heat exchangers[6-8]. In particular, open-cell metallic foams exhibit high fluid permeability and a large
specific surface area thanks to their high porosity, making them suitable for energy
absorption and filtration applications[9-12]. As research in environmental and alternative energy fields increases, interest in
the potential application of metallic foam as a catalyst matrix are growing[11,12].
Catalysts are inevitably exposed to high-temperature environments to maintain reaction
efficiency. So metallic foams intended for use as catalyst matrices must exhibit excellent
thermal and high-temperature mechanical properties. Furthermore, depending on the
reactor environment, their energy absorption properties at both room and elevated
temperatures must be considered. To fabricate Ni superalloy metallic foams, various
processes—such as electron-beam directed vapor deposition[13], casting replication[14], slurry foaming[15], combustion synthesis[16], powder sintering[17], spark plasma sintering[18], and pack-cementation[19-22]—have been attempted. But these methods faced practical challenges, including difficulties
in controlling cell size and the distribution of alloying elements, as well as high
processing costs, process complexity, and potential of contamination during process.
And as noted earlier, to obtain the desirable mechanical properties and oxidation
resistance of Ni-based superalloys requires precise control of alloy composition.
However, achieving simultaneous control of alloy composition and foam morphology remains
challenging in these methods, as alloying is predominantly governed by diffusion processes,
which makes precise control over the distribution of alloying elements inherently
difficult, while also presenting practical challenges in controlling the geometry
of the porous structure, thereby hindering independent control of these two critical
factors.
To overcome these limitations, a study has attempted to fabricate Ni-superalloy foams
using a powder spraying process[23,24]. This process involves spraying and adsorbing alloying powders onto pure Ni foam,
followed by debinding and sintering steps. It offers the advantage of easy control
over metal composition by varying powder size, quantity, and alloying elements. Furthermore,
the shape and surface area of the porous structure can be tailored based on the polyurethane
pre-form. In short, because the powder spraying/sintering process consists of two
distinct stages—urethane pre-form shaping followed by powder deposition and sintering—it
enables independent yet simultaneous control of morphology and composition of the
foam through tailoring of the pre-form and regulation of the size, composition, and
distribution of the sprayed powders. Previous studies on Ni-Fe-Cr-Al superalloy foams
produced via this method have investigated their surface area, room-temperature compressive
strength[23], hot corrosion properties[25], and creep properties[26]. However, the mechanical properties of metallic foams must be evaluated by considering
both their structural characteristics and the influence of microstructure according
to alloy composition, yet research in this area remains insufficient. In addition,
to expand the applications of Ni-based superalloy foams to fields such as impact absorbers
as well as catalyst matrices, their energy absorption behavior at both room and elevated
temperatures is a critical consideration. Nevertheless, there is a lack of research
that analyzes this behavior in connection with the microstructure.
For these reasons, the room- and high-temperature compressive properties of Ni–Fe–Cr–Al
alloy foams fabricated via a powder spraying process were systematically compared
as a function of pore size in this study, with simultaneous consideration of microstructural
features. Since pore size is closely associated with the relative density and geometry
of the foam, it significantly influences key mechanical properties, such as plateau
strength and compressive deformation behavior, thereby determining the energy absorption
characteristics of the foam. In addition, as the powder spraying/sintering process
enables precise control over microstructure and considering that microstructural features
can significantly affect the deformation behavior of the foam, microstructural features
such as phase distribution and the formation of brittle phases, were also analyzed
and correlated with the compressive deformation behavior and energy absorption characteristics
of the foam.
2. EXPERIMENTAL METHODS
To fabricate Ni–Fe–Cr–Al alloy foams, a binder and Fe–Cr–Al–based alloy powders (35.56
wt.% Fe, 34.92 wt.% Cr, 20.0 wt.% Ni, and 9.52 wt.% Al) were sprayed and distributed
onto Ni pre-foams. The powder-to-foam ratio (PFR), defined as the weight percentage
of powder relative to the total mass of the alloyed foam, was controlled at 62% for
the 800 μm pore foam and 63% for the 580 μm pore foam. Subsequently, Ni–22.4 wt.%
Fe–22.0 wt.% Cr–6.0 wt.% Al alloy foams were successfully fabricated through debinding
and transient liquid phase sintering heat treatments.
The structure and surface morphology of the fabricated Ni-based alloy foams were examined
using scanning electron microscopy (SEM; TESCAN, VEGA II). An image analyzer was employed
to quantify the average pore size, strut thickness, and wall thickness. Phase analysis
and microstructural characterization of the Ni–Fe–Cr–Al alloy foams were conducted
using X-ray diffraction (XRD; RIGAKU D2000) and standard metallographic preparation
techniques. The specimens were ground using SiC emery papers and subsequently mirror-polished
with a 1 μm Al2O3 slurry. Etching was performed for several seconds using a solution of 25 mL HCl,
25 mL H2O, and 10 mL H2O2, after which the microstructures were observed by SEM. The specific surface areas
of the samples were measured using a specific surface area analyzer (3Flex, Micromeritics)
based on BET method.
To evaluate the room and high temperature compressive properties of the Ni–Fe–Cr–Al
alloy foams, specimens were prepared with dimensions of 15 mm (L) × 5 mm (W). The
thicknesses were set to 2.5 mm for the 800 μm foam and 1.9 mm for the 580 μm foam.
Compression tests were carried out using an INSTRON 8801 universal testing machine
at an initial strain rate of 1×10-3 s-1. For high-temperature testing, a halogen furnace was employed at a heating rate of
20 K/min. Specimens were held at the target temperatures (873 K, 973 K, and 1073 K)
for 10 minutes to ensure thermal stabilization. Subsequently, tests were performed
at the same initial strain rate until a final compressive strain of 60% was reached.
After compression tests, the deformed morphologies were examined via SEM to analyze
the temperature-dependent compressive deformation, energy absorption, and fracture
behaviors. For each temperature condition, three independent compression tests were
conducted to ensure reproducibility and the average values with corresponding variations
were used for analysis.
3. RESULTS and DISCUSSION
3.1 Structural and microstructural characteristics of Ni-Fe-Cr-Al superalloy foams
Figure 1 presents the representative architecture of the Ni–Fe–Cr–Al alloy foams fabricated
via the powder spraying/sintering process. The foams exhibit an open-pore, three-dimensional
reticulated structure composed of cells with a tetrakaidecahedral geometry, each consisting
of approximately 12–14 pentagonal or hexagonal faces, irrespective of the average
pore size. The average pore sizes were 787 μm (hereafter denoted as the 800 foam)
and 562 μm (hereafter denoted as the 580 foam), with corresponding average strut thicknesses
of 97.2 μm and 58.6 μm, respectively. The struts exhibited a hollow triangular pipe-like
morphology, originating from the polyurethane (PU) preform employed during fabrication.
The cell walls forming the struts consisted of grains with sizes on the order of 10–20
μm, and the average wall thicknesses were measured to be 15.5 μm for the 800 foam
and 12.6 μm for the 580 foam. The relative densities were determined to be 6.19% and
8.51% for the 800 and 580 foams, respectively, and the detailed structural parameters
are summarized in Table 1.
Fig. 1. SEM micrographs of Ni-Fe-Cr-Al superalloy foams; (a) 800 foam and (b) 580
foam
Table 1. Geometric parameters of the Ni-Fe-Cr-Al superalloy foams
|
Sample name
|
Relative
Density (%)
|
Apparent
Density (g/cm3)
|
Pore
Size (μm)
|
Strut
Thickness (μm)
|
Cell wall
Thickness (μm)
|
|
Ni-Fe-Cr-Al foams
|
800
|
6.19
|
0.49
|
787
|
97.2
|
15.5
|
|
580
|
8.51
|
0.68
|
562
|
58.6
|
12.6
|
The specific surface areas were measured as 0.0149 ± 0.0042 m2/g for the 800 foam and 0.0580 ± 0.0070 m2/g for the 580 foam. The markedly higher surface area of the 580 foam is attributed
to the rough strut surface morphology induced by adhered powder particles, as well
as its smaller pore size, which increases the strut density per unit volume and consequently
provides a larger effective surface area for powder attachment.
Figure 2 shows the surface morphology, cross-sectional features, and microstructural characteristics
of the struts. The strut surfaces are covered with powders adhered during the powder
spraying process used to fabricate the Ni-based superalloy foams (Fig. 2(a)). Cross-sectional observations reveal sound metallurgical bonding between the strut
and the deposited powders, with no evidence of oxide layers, interfacial detachment,
or cracking (Fig. 2(b)). This bonding behavior was consistently observed in both foams, irrespective of
pore size.
Fig. 2. SEM micrographs of typical microstructures of Ni-Fe-Cr-Al superalloy foams
(800 foam); (a) surface morphologies of the strut, (b) cross-section and (c) grain
sizes in the strut
Compared with Ni-based foams fabricated by the pack-cementation process[19-22], the present foams exhibit a more homogeneous distribution of alloying elements (Fig. 3). This difference results from the distinct characteristics of the powder spraying/sintering
and pack-cementation processes. In the pack-cementation method, alloying is governed
by diffusion, and the non-uniform distribution of packing materials within the porous
structure can lead to variations in diffusion behavior, resulting in compositional
inhomogeneity. In contrast, in the powder spraying/sintering process, powders are
more uniformly deposited on the strut surfaces, and the alloy composition is largely
determined by the amount and distribution of sprayed powder prior to sintering. This
enables more precise control of both alloy composition and microstructure within the
porous structure.
Fig. 3. Distribution of alloying elements of as-fabricated Ni-Fe-Cr-Al superalloy
580 foam;
Al-rich regions were observed within the strut microstructure, originating from the
constituent phases of the Ni-based superalloy. To identify these phases, XRD and EDS
analyses were conducted, as shown in Fig. 4. The XRD patterns (Fig. 4(a)) indicate that both foams, despite their different pore sizes, consist of γ-Ni (JCPDS
#04-0850), γ′-Ni (JCPDS #04-0850), and β-NiAl (JCPDS #44-1185) phases[24,25]. Based on a combined analysis of the XRD and EDS results, the Al-rich regions observed
in Fig. 4(b) were identified as the β-NiAl phase (indicated by black arrows).
Fig. 4. (a) X-ray diffraction analysis results of Ni-Fe-Cr-Al foams (closed square:
γ-Ni, open square: γ’, and open circle: β-NiAl) and (b) EDS results
3.2 Compressive deformation behavior of Ni-Fe-Cr-Al superalloy foams from RT to High
Temperature
Figure 5 illustrates the compressive deformation behavior of the Ni–Fe–Cr–Al alloy foams at
room and elevated temperatures, categorized according to their initial pore sizes.
The Ni–Fe–Cr–Al superalloy foams exhibited the three characteristic stages of cellular
material compression: (a) the linear elastic stage, (b) the plastic collapse stage,
and (c) the densification stage. The compressive properties of the 800 and 580 foams
at room and high temperatures are summarized in Table 2.
Fig. 5. The compressive stress-strain curves of the Ni-Fe-Cr-Al superalloy foams from
room to 1073 K; (a) 800 foam and (b) 580 foam
Table 2. The mechanical properties of the Ni-Fe-Cr-Al superalloy foams from RT to
high temperature (σₗ: plateau strength, σp: average stress over the plateau region)
|
Sample name
|
Temperature (K)
|
ρ*/ρ (%)
|
σₗ(MPa)
|
σP (MPa)
|
ΔεP (%)
|
|
800 foam
|
RT
|
6.19
|
2.98
|
2.92
|
44
|
|
873 K
|
2.70
|
2.84
|
47
|
|
973 K
|
2.28
|
1.97
|
48
|
|
1073 K
|
1.55
|
1.12
|
48
|
|
580 foam
|
RT
|
8.51
|
4.38
|
4.40
|
38
|
|
873 K
|
4.17
|
4.22
|
42
|
|
973 K
|
2.85
|
2.71
|
46
|
|
1073 K
|
1.42
|
1.42
|
46
|
At room temperature, the plateau strength of the 800 foam was measured to be 2.98
MPa, whereas the 580 foam exhibited a higher plateau strength of 4.38 MPa.
Notably, the plateau strengths of the Ni–Fe–Cr–Al alloy foams obtained in this study
were higher than those reported for Ni-based foams with similar relative densities
in previous studies[19-22]. This improved performance is attributed to the more uniform elemental distribution
and the effective control of alloying element concentrations achieved by the present
fabrication process. Furthermore, the flow curves of the Ni–Fe–Cr–Al alloy foams do
not exhibit a significant stress drop following the initial peak stress. Compared
with the results reported by H. Choe and D. C. Dunand[19], the foams in the present study showed relatively smooth flow stress curves. This
behavior can be associated with the relatively small pore size and the distribution
of the β-NiAl phase within the Ni matrix. Especially the absence of a pronounced stress
drop suggests that relatively uniform phase distribution may help to mitigate localized
stress concentration, thereby promoting stable plastic deformation which leads to
favorable energy absorption capability. This behavior is discussed in further detail
in Section 3.3.
In the high-temperature compression tests, the alloy foams showed the same three stages
of compressive deformation observed at room temperature: elastic deformation, plastic
collapse, and densification. Overall, the compressive strength decreased monotonically
with increasing temperature. Notably, irrespective of pore size, the compressive stress–strain
curves obtained at 873 K were remarkably similar to those measured at room temperature.
For the 800 foam, the plateau strength decreased from 2.98 MPa at room temperature
(RT) to 2.70 MPa at 873 K, 2.28 MPa at 973 K, and 1.55 MPa at 1073 K. In contrast,
the 580 foam exhibited plateau strengths of 4.38 MPa (RT), 4.17 MPa (873 K), 2.85
MPa (973 K), and 1.42 MPa at 1073 K. These results indicate that, although both foams
exhibit comparable plateau strengths up to 873 K, a noticeable decrease occurs at
higher temperatures (973 K and 1073 K). When evaluated relative to 873 K, the average
plateau stress decreases by approximately 30.6% and 60.6% for the 800 foam, and 35.8%
and 66.4% for the 580 foam at 973 K and 1073 K, respectively. These results demonstrate
substantial strength degradation at elevated temperatures rather than minor variations.
In general, the plateau strength of cellular materials is strongly correlated with
their relative density according to the Gibson–Ashby model[6], which can be expressed as
In this relationship, $\sigma^{*}_{pl,foam}$ denotes the plastic-collapse strength
of the foam, σys represents the yield strength of the bulk material, ρ* is the apparent density of the foam, and ρs is the density of the fully dense bulk material. In this study, ρs represents the theoretical density and was calculated from the alloy composition
using the rule of mixtures. Accordingly, ρ*/ ρs corresponds to the relative density. The constant C is a geometric factor, typically
ranging from 0.23 to 0.30. Based on this relationship, the bulk yield strength of
the material constituting the Ni–Fe–Cr–Al foam was estimated as a function of relative
density. the results are presented in Figure 6.
Fig. 6. (a) Temperature effects on the plateau strengths of Ni-Fe-Cr-Al foams and
(b) estimated bulk yield strengths of Ni-based superalloy foam compared with reported
strengths.
The calculated bulk yield strengths for the 800 foam were 643.4 MPa at room temperature
(RT), 583.0 MPa at 873 K, 492.3 MPa at 973 K, and 334.7 MPa at 1073 K. For the 580
foam, the corresponding values were 586.0 MPa (RT), 557.9 MPa (873 K), 381.3 MPa (973
K), and 190.0 MPa (1073 K), demonstrating a clear difference between the two foams
with different pore sizes.
According to J. Huang et al.[5], wrought Ni–Fe–Cr–Al superalloys exhibit yield strengths of approximately 410 MPa
at 973 K and 270 MPa at 1073 K. Compared with these values, the estimated bulk yield
strengths of the 800 foam are slightly higher, whereas those of the 580 foam are relatively
lower. This discrepancy is considered to arise primarily from pore-size-dependent
structural factors, including variations in node density (which may induce constraint
effects during deformation), wall thickness, and strut thickness. In addition, in
the present study, microstructural features, such as the distribution of brittle phases,
were also considered in interpreting the mechanical response of the foams.
3.3 Energy-Absorption behaviors of Ni-Fe-Cr-Al superalloy foams from RT to High Temperature
One of the most important mechanical properties of cellular materials is their energy
absorption capability. In particular, for Ni–Fe–Cr–Al superalloy foams, it is essential
to analyze the high-temperature energy absorption behavior under quasi-static or dynamic
loading conditions. Among the parameters governing the energy absorption performance
of foams, the flow stress and densification strain are the most critical.
The plateau regime corresponds to the stage in which strut bending and compression
occur progressively, resulting in either a nearly constant flow stress or a slight
increase in flow stress with increasing strain on the stress–strain curve. In the
densification regime, the stress rises sharply due to interaction (mechanical anchoring
or friction) between bending/compression struts. However, these two deformation behaviors
do not occur sequentially; instead, they develop simultaneously with increasing strain,
making it difficult to clearly distinguish between the plateau and densification regimes,
as shown in Figure 7.
Fig. 7. Typical deformation behavior of Ni-Fe-Cr-Al superalloy foams with respect
to increasing compressive strains at room temperature (bending and collapsed struts
are indicated by black arrows)
Therefore, in this study, the evolution of energy absorption was continuously evaluated
as a function of compressive strain. The amount of energy absorbed with increasing
compressive strain was calculated using the following equation[6].
In the equation, W represents the energy absorbed by the cellular material, and ε
denotes the compressive strain.
Figure 8 shows the energy absorption behavior of the Ni–Fe–Cr–Al alloy foams as a function
of compressive strain. The 800 foam exhibited a lower energy absorption capacity than
the 580 foam, which is associated with its lower strength resulting from the lower
relative density.
Fig. 8. Absorbed energy per unit volume for deformation to increasing compressive
strain; (a) 800 foam and (b) 580 foam from room temperature to high temperatures
A characteristic feature observed in the energy absorption curves is that, irrespective
of pore size and temperature, the absorbed energy increases approximately linearly
up to a certain compressive strain, followed by a region in which the energy absorption
rate increases rapidly. This change in slope is associated with the onset of densification,
during which a rapid increase in strength occurs. With increasing temperature, the
compressive strain corresponding to this slope change shifts to higher values, suggesting
a delayed densification behavior at elevated temperatures. This tendency may be related
to the enhanced ductility of the Ni–Fe–Cr–Al alloy at high temperatures. In addition,
the 800 foam exhibited a higher densification strain than the 580 foam over the entire
temperature range examined. This behavior is consistent with previously reported trends
indicating that the densification strain generally increases with decreasing relative
density[6-7].
To compare the energy absorption capacity of the Ni–Fe–Cr–Al alloy foam investigated
in this study, its performance was evaluated against that of a Ni-based alloy foam
with a similar relative density reported by Q. Pang et al.[20], as shown in Figure 9.
Fig. 9. Relative density vs. absorbed energy of the Ni-Fe-Cr-Al foams developed in
this study from room temperature (RT) to 1073 K. The energy absorption at RT is compared
with those of commercially pure Ni foam and the Ni-Fe-Cr foams (before and after homogenization)
reported by Q. Pang et al.
The results indicate that the Ni–Fe–Cr–Al alloy foam exhibits higher energy absorption
than the referenced Ni-based alloy foam, and the absorbed energy values for each data
point are as follows : for the 580 foam, the values are 2.72 at RT, 2.74 at 873 K,
1.57 at 973 K, and 0.91 J/cm2 at 1073 K, while for the 800 foam, the corresponding values are 1.52 at RT, 1.55
at 873 K, 1.13 at 973 K, and 0.64 J/cm2 at 1073 K. according to these results, the energy absorption is nearly maintained
up to 873 K, irrespective of the pore size.
In general, increasing temperature leads to a reduction in strength and an increase
in ductility for metallic materials. However, for Ni-based superalloys, it has been
reported that within the temperature range of approximately 873–973 K, changes in
the operative slip systems of the γ′ phase result in anomalous yield strength behavior,
commonly referred to as super-heat-resistant behavior[1]. Furthermore, previous studies have shown that the yield stress of the β-NiAl phase
is maintained up to around 800 K, followed by a rapid decrease at temperatures above
approximately 900 K[27,28]
Based on the microstructural characteristics of the Ni–Fe–Cr–Al alloy foam examined
in this study, the retention of strength and energy absorption at 873 K can be associated
with the super-heat-resistant behavior of the γ′ phase together with the relatively
stable yield strength of the β-NiAl phase (Figures 6 and 8). At temperatures above this range, softening of both the β-NiAl phase and the matrix
is expected to occur, leading to a pronounced reduction in strength and, consequently,
a decrease in energy absorption capacity.
With respect to pore size, although the 580 foam exhibits a lower densification strain
(εd) than the 800 foam, it consistently absorbs a greater amount of energy. Moreover,
at the same compressive strain, the energy absorbed by the 580 foam is higher than
that of the 800 foam. This behavior is associated with the higher plateau strength
and lower densification strain of the 580 foam, which promote earlier densification
and a more rapid increase in flow stress.
3.4 Effects of constituent phases on the mechanical properties of Ni-Fe-Cr-Al foams
The compressive deformation modes were examined to elucidate the influence of microstructure
on compressive deformation behavior and energy absorption, and the corresponding results
are shown in Figure 10.
Fig. 10. Fractured surface of bent and collapsed struts in the Ni-Fe-Cr-Al foams with
respect to increasing temperatures; at (a) RT, (b) 873 K and (c) 1073 K
At room temperature, bending-dominated deformation of the struts was predominantly
observed in the Ni–Fe–Cr–Al alloy foams, irrespective of pore size. In struts subjected
to severe bending deformation, fracture surfaces were frequently detected. At room
temperature, mixed-mode fracture surfaces consisting of cleavage and dimpled features
were observed for both foams, independent of pore size.
In the case of cleavage fracture, as shown in Figure 11 (a), fracture was observed to occur primarily in regions exhibiting a relatively dark
gray contrast, which correspond to the β-NiAl phase. In addition, crack initiation
at the interface between the β-NiAl phase and the surrounding matrix was frequently
observed, as indicated by the white arrow.
Fig. 11. Effects of constituent phases on the fracture behavior of struts in Ni-Fe-Cr-Al
foams at (a) room temperature and (b) 1073 K
Detailed observations of the specimens before and after deformation revealed that
the β-NiAl phase is predominantly formed at the interface between the pure Ni strut
and deposited powders, as well as along the grain boundaries of the Ni strut, where
fast diffusion occurs, as shown in Figures 3 and 4. During deformation, the foam structure accommodates strain through bending and compression
of the struts (Fig. 7). In particular, tensile stress generated on the strut surface during bending leads
to preferential crack initiation in the brittle β-NiAl phase located near the surface,
followed by crack propagation into the interior, which is consistent with the fracture
features observed in Figure 11 (a).
At elevated temperatures, the fracture morphology gradually transitions to dimpled
or ductile fracture features associated with plastic deformation. This indicates that
ductile fracture becomes dominant at elevated temperatures, whereas brittle fracture
associated with the β-NiAl phase is predominant at room temperature. In contrast,
ductile fracture accompanied by dimple formation and plastic deformation was primarily
observed in the γ-Ni matrix, as shown in Figure 10(b–c), indicated by white arrows.
In Ni-based superalloys, precise control of chemical composition is a critical factor
governing hot corrosion resistance and mechanical properties. In particular, increasing
the Al content has been shown to significantly improve hot corrosion resistance through
the formation of a protective Al2O3 scale[3-4]. At the same time, Al plays a key role in the formation of strengthening phases such
as γ′-Ni3Al and β-NiAl, and thus can markedly influence the mechanical response of the alloy[29,30]. With increasing Al content, the formation of brittle phases, including β-NiAl, becomes
more pronounced, which may lead to a reduction in material toughness.
This effect is especially important in metallic foams with thin struts, where changes
in ductility and brittleness can strongly affect deformation and fracture behavior.
Accordingly, control of alloy composition can be considered an important processing
parameter for tailoring the plateau strength and energy absorption characteristics
of Ni-based superalloy foams, particularly through its influence on the formation
and distribution of brittle phases such as β-NiAl. The results of the present study
suggest that the formation of foams exhibiting a balanced combination of strength,
plastically deformable phases, and thermally stable phases can be achieved through
the powder spraying/sintering process.
4. CONCLUSIONS
In this study, the room- and high-temperature compressive deformation behaviors and
energy absorption characteristics of Ni–Fe–Cr–Al superalloy foams fabricated via powder
spraying and subsequent sintering were systematically investigated as a function of
pore size. The main conclusions can be summarized as follows :
Alloying powders were successfully adhered to the struts of the foams fabricated via
the powder spraying/sintering process. The average pore sizes were measured to be
787 μm for the 800 foam and 562 μm for the 580 foam. The corresponding average strut
thicknesses were 97.2 μm and 58.6 μm, respectively, while the relative densities were
determined to be 6.19% for the 800 foam and 8.51% for the 580 foam.
Irrespective of pore size, the Ni–Fe–Cr–Al superalloy foams were composed of a γ matrix
together with γ′ and β-NiAl phases. The γ matrix forming the cell walls exhibited
a grain size of approximately 10~20 μm. Energy-dispersive spectroscopy (EDS) analysis
results indicated that the β-NiAl phase was distributed non-uniformly within the struts.
The plateau strength of the superalloy foams showed a clear dependence on pore size.
At room temperature, the plateau strength of the 580 foam (4.38 MPa) was higher than
that of the 800 foam (2.98 MPa), exceeding those of Ni-based foams with comparable
relative densities reported in the literature. The plateau strength was maintained
up to 873 K, whereas at higher temperatures (~1073 K) it decreased to 1.42 MPa for
the 580 foam and 1.55 MPa for the 800 foam. As a result, the energy absorption capacity
was preserved up to 873 K, indicating favorable high-temperature energy absorption
behavior. This behavior is associated with the anomalous strengthening of the γ′ phase
and the thermal stability of the β-NiAl phase in the Ni–Fe–Cr–Al superalloy foam.
At room temperature, the Ni–Fe–Cr–Al superalloy foams exhibited cleavage-type fracture
associated with grain boundary fracture (intergranular fracture) and brittle fracture
within the β-NiAl phase. Similar fracture characteristics were observed up to 873
K, where the mechanical properties were retained. At temperatures above this range,
the fracture mode transitioned to ductile fracture accompanied by plastic deformation.