Gamma Prime Precipitate Evolution During Aging of a Model Nickel-Based Superalloy
Gamma Prime Precipitate Evolution During Aging of a Model Nickel-Based Superalloy
A.J. GOODFELLOW 0
E.I. GALINDO-NAVA 0
K.A. CHRISTOFIDOU 0
N.G. JONES 0
T. MARTIN 0
P.A.J. BAGOT 0
C.D. BOYER 0
M.C. HARDY 0
H.J. STONE 0
0 A.J. GOODFELLOW, E.I. GALINDO-NAVA, K.A. CHRISTOFIDOU, N.G. JONES, and H.J. STONE are with the Department of Materials Science and Metallurgy, University of Cambridge , 27 Charles Babbage Road, Cambridge CB3 0FS , UK. Contact
The microstructural stability of nickel-based superalloys is critical for maintaining alloy performance during service in gas turbine engines. In this study, the precipitate evolution in a model polycrystalline Ni-based superalloy during aging to 1000 hours has been studied via transmission electron microscopy, atom probe tomography, and neutron diffraction. Variations in phase composition and precipitate morphology, size, and volume fraction were observed during aging, while the constrained lattice misfit remained constant at approximately zero. The experimental composition of the γ matrix phase was consistent with thermodynamic equilibrium predictions, while significant differences were identified between the experimental and predicted results from the γ′ phase. These results have implications for the evolution of mechanical properties in service and their prediction using modeling methods.
POLYCRYSTALLINE Ni-based superalloys are the
material of choice for many high-temperature structural
applications in gas turbine engines. Their remarkable
mechanical performance is derived from the presence of
an ordered L12 (strukturbericht notation) γ′ precipitate
phase within the disordered A1 γ matrix. The principal
mechanisms by which these alloys are strengthened
include order and coherency strengthening from the γ′
precipitates, as well as solid solution strengthening of
the γ matrix phase and grain boundary hardening.[
the extent of precipitation strengthening is dependent on
the γ′ morphology and particle size distribution, these
are carefully controlled through sagacious selection of
heat treatments. However, care is taken to ensure that
the resultant microstructure delivers an appropriate
balance of properties.
The precipitate size at which peak strength is obtained
corresponds to the transition from weak to strong pair
dislocation coupling and, as such, varies with alloy
composition. For example, this occurs for 55 to 85 nm γ′
precipitates in Nimonic 105, but just 26 to 30 nm for
] However, a unimodal distribution of such fine
precipitates is difficult to achieve in practice and has
been associated with reduced ductility[
] and creep
] As such, a multimodal particle size
distribution is often utilized, comprising secondary
(∼100 to 250 nm) and tertiary (∼5 to 50 nm)
intragranular precipitates. Larger, primary (∼1 µm) intergranular
precipitates may also be present, depending upon the
solution heat treatment temperature.
The evolution of γ′ particle size distributions during
high-temperature service will affect the mechanical
properties and must therefore be understood.
Computational tools based upon thermodynamic and kinetic
models exist to predict such precipitate evolution.[
However, further experimental studies are needed to
validate the assumptions made in these models,
particularly the compositions of the γ′ precipitates, as these are
known to differ between the primary, secondary, and
tertiary distributions and will change with time at
In this study, the γ and γ′ phase compositions and the
distribution of the γ′ precipitates have been investigated
in a model superalloy as a function of time at
temperature using a combination of transmission electron
microscopy and atom probe tomography. The resulting
effect on lattice misfit has also been studied using
neutron diffraction. These data identify the key
discrepancies between thermodynamically predicted and
experimentally measured precipitate compositions, which
may have significant implications for the modeling of
precipitate evolution and alloy strength.
II. EXPERIMENTAL DETAILS
To characterize the effect of aging on microstructure,
a model Ni-based superalloy was designed, with
nominal composition given in Table I. This composition lies
within the range of current, high-strength polycrystalline
Ni-based superalloys, such as RR1000,[
] and is
intended to have low lattice misfit between the γ matrix
and γ′ precipitate phases at room temperature,
comparable to certain commercial disk alloys.[
] The alloy
was manufactured through vacuum induction melting
(VIM) using elements of ≥ 99.9 pct. purity and cast into
10 mm-diameter cylindrical bars using steel molds.
Differential Scanning Calorimetry (DSC) was carried
out on the as-cast material using a NETZSCH 404
instrument. Samples were heated at 10 K min−1 to 1723 K
(1450 °C), before cooling at the same rate to room
temperature. From the thermograms acquired, the γ′
solvus and solidus temperatures were determined to be
1448 K and 1560 K (1175 °C and 1287 °C), respectively.
This enabled a suitable homogenization heat treatment
temperature to be identified within the single γ phase field.
Homogenization of the alloy was conducted for 22 hours
at 1523 K (1250 °C) in an Ar-backfilled quartz ampoule, to
avoid oxidation. Subsequent air cooling occurred at
an average rate of ∼7.8 K s−1 for the initial 200 K, slowing
to an average of ∼1.6 K s−1 between 1273 K and 673 K
(1000 °C and 400 °C). Sections of the homogenized alloy
were subjected to aging heat treatments at 1033 K (760 °C)
for durations of 1, 2, 16, 100, 200, and 1000 hours, followed
by air cooling at an average rate of ∼4.4 K s−1
Thermodynamic modeling to determine the
equilibrium composition of the matrix and γ′ precipitate phases
at 1033 K (760 °C) was conducted using the
ThermoCalc software package with the TCNi7 database,
excluding all other phases. The lattice parameters of the
γ and γ′ phases (aγ and aγ′ respectively) were obtained
from the predicted molar volumes (Vm) using Eq. [
and the lattice misfit (δ) was then calculated from
ðac0 þ acÞ
Electron transparent samples of the heat-treated alloy
for microstructural analysis in a transmission electron
microscope (TEM) were prepared by electropolishing
and through the production of carbon replicas.
Electropolishing was performed in a 5 pct. perchloric acid in
methanol solution at a temperature of 268 K (−5 °C)
and a voltage of 20 V.
Carbon replica samples were obtained by
electrolytically etching material polished to a 0.5 µm finish in a 10
pct. phosphoric acid in water solution at ∼3 V until a
blue halo appeared. After washing in ethanol, formvar
in chloroform solution was deposited on the sample
surface and used to attach a piece of acetate sheet. Once
dry, the acetate sheet was removed to eliminate any
over-etched γ′ particles from the surface. The samples
were then sputter coated with carbon, and this carbon
coating was scored into ∼2 mm squares. Finally,
electrolytic etching in a 20 pct. perchloric acid in ethanol
solution at 10 V was carried out until the carbon coat
began to blister. These fragments were then floated onto
copper TEM grids for analysis.
Microstructural imaging was carried out by Scanning
Transmission Electron Microscopy (STEM) using an FEI
Tecnai Osiris TEM, operated at an accelerating voltage of
200 keV. Compositional analyses were performed by
Energy-Dispersive X-ray Spectroscopy (EDX) in the same
instrument, using an FEI Super-X EDX detector.
The electropolished samples were used for
compositional analysis of the γ phase using STEM–EDX. Data
were acquired from regions of these samples between the
γ′ precipitates. The same method could not reliably be
used for the γ′ precipitates, as it was not possible to
ensure that the electron beam did not pass through γ
phase above or below the γ′ precipitates. To address this,
compositional analysis of the γ′ precipitates was
performed on the samples that were prepared as carbon
extraction replicas, which were devoid of the γ phase.
Particle size distributions (PSDs) for each heat treatment
were determined by manually tracing around a minimum of
300 secondary γ′ particles on images of the carbon replica
samples. This approach was required due to the presence of
overlapping particles, which cannot be separated
automatically. Analysis of the precipitate outlines was performed
] from which the equivalent circular
diameters were obtained. Tertiary γ′ analysis was carried out via
the same method but using higher resolution STEM
micrographs. The equivalent circular diameters were binned
using the Freedman–Diaconis method[
] and fitted with
lognormal functions, given in Eq. [
], using Igor Pro
(WaveMetrics, Lake Oswego, OR). The average precipitate
size after each heat treatment was taken as the median of
the lognormal functions (eμ) and the associated uncertainties
were taken to be the width of the function away from the
median value (e(μ+σ) and e(μ−σ)), where σ. is the error
(standard deviation) associated with the coefficient given by
the Igor Pro software.
fðxÞ ¼ xwpffi2ffiffipffiffi exp
To determine the volume fractions of the secondary γ′
precipitates, Scanning Electron Microscopy (SEM) was
performed using an FEI Nova NanoSEM FEG SEM
operated at ∼5 keV. Samples in each heat-treated
condition were polished to 0.06 µm finish with colloidal
silica and imaged in Backscattered Electron Imaging
(BSE) mode. Quantification of the secondary γ′ volume
fraction was achieved using thresholded images in
ImageJ. The uncertainty quoted with the volume
fraction measurements was taken to be the standard
deviation of the volume fraction in four separate areas
across each sample.
Atom Probe Tomography (APT) was performed for
compositional analysis of the constituent phases and
comparison with the results obtained using STEM–
EDX. APT samples were prepared from needles with a
square cross-section of length 0.5 mm. Subsequent
electropolishing was performed in two stages: firstly using a
solution of 10 pct. perchloric acid in acetic acid at a
voltage of 22 V, and secondly using a solution of 2 pct.
perchloric acid in 2-butoxyethanol at 23 V. Data were
acquired from each electropolished needle using a LEAP
5000 XR instrument in laser mode with a wavelength of
355 nm. APT data were acquired at 50 Κ (−223 °C), a
pulse rate of 200 kHz, and a pulse energy of 50 pJ.
Neutron diffraction was carried out on the C2 powder
diffractometer at the Canadian Neutron Beam Centre
(CNBC), Chalk River, Canada. The wavelength of the
incident beam was determined to be ∼1.33 A˚ using an
Al2O3 standard. Data were acquired from each sample for
3 hours at room temperature, using a position-sensitive
detector over a 2θ range of 36 to 116 deg. The samples
were rotated during data collection to improve the
counting statistics. For each sample, the γ′ superlattice
reflections were fitted individually in WaveMetrics Igor
Pro using Gaussian functions. The associated lattice
parameter was determined from the fitted superlattice
positions using the Nelson–Riley method, in which the
equivalent lattice parameters from each peak were plotted
against the associated absorption error, (δa/a)abs, Eq. [
and the lattice parameter was taken as the linear
] The positions of the γ peaks were obtained by
fitting the fundamental reflections with two Gaussian
functions, one of which was constrained to the position
associated with the previously determined γ′ lattice
parameter. As before, the γ lattice parameter was obtained from
the fitted peak positions using the Nelson–Riley function,
]. Conversion of the measured lattice parameters to
lattice misfit was performed using Eq. [
1 cos2 h
ðda=aÞabs ¼ 2
sin h þ
STEM images of the γ′ phase from the carbon replicas
of the alloy in each heat treatment condition are shown
in Figure 1. With the exception of the sample aged for
1000 hours, all conditions exhibited bimodal PSDs of
secondary and tertiary γ′. As a consequence of the
different precipitate sizes, images of the secondary
(Figure 1, left-hand side images) and tertiary γ′ (Figure 1,
right-hand side images) have been displayed at different
The morphology of the secondary γ′ showed no
appreciable change with aging time up to 16 hours, with
all secondary precipitates remaining approximately
cuboidal. On further aging to 100 hours, octodendritic
precipitates alongside smaller, cuboidal precipitates
were observed, characteristic of the morphological
instabilities associated with precipitate splitting.[
Following aging for 200 hours, only octodendritic
precipitates were observed throughout the sample. After
1000 hours, the secondary γ′ appeared finer and more
spherical than those observed during shorter duration
exposures. In contrast to the secondary γ′, the tertiary γ′
remained approximately spherical throughout aging and
increased continuously in size up to 200 hours, after
which they were no longer present within the
PSDs of both the secondary and tertiary γ′ obtained
from analysis of the STEM images of the
as-homogenized condition and following aging for 1, 16, 100, and
1000 hours are shown in Figure 2. All PSDs were well
described by a lognormal function. The single additional
peak visible between the distributions of secondary and
tertiary γ′ precipitates in the alloy after aging for 16
hours is due to the presence of one large tertiary
precipitate and one small secondary precipitate, of
similar intermediate size being identified in the
microstructural regions examined. Example PSDs of
the secondary and tertiary precipitates, alongside the
fitted lognormal distributions, for the sample aged for
16 hours are provided in the Supplementary
Information to demonstrate the quality of the fit to the data.
The median diameters of the secondary and tertiary γ′
precipitates are presented in Figure 3. The median
diameters of the secondary γ′ remained approximately
constant, ∼200 nm, throughout aging, although the
sample aged for 100 hours displayed a noticeably lower
median value. This was also associated with a broader
PSD, as indicated by the larger error bars quoted with
this data point. This arose as a result of the presence of
regions containing smaller, cuboidal precipitates
alongside regions of larger, octodendritic precipitates in this
sample. With increasing aging time, the size of the
tertiary γ′ increased monotonically, accompanied by a
progressive increase in the width of the PSDs. Unlike the
secondary γ′, no changes associated with splitting
instabilities were observed.
The volume fractions of secondary precipitates
obtained from analysis of the BSE images are presented
in Figure 4. These results indicate that the volume
fraction increases with aging time from 37(±1) pct. after
1 hour to 46(±3) pct. after 1000 hours. Given the
constant precipitate size shown in Figure 3, this suggests
that the number density of secondary γ′ precipitates
increases during aging. The larger uncertainty associated
with the volume fraction measurements after 100 hours
may be attributed to variations in the observed volume
fractions in regions of the sample at different stages of
precipitate splitting. The volume fraction of the tertiary
γ′ was observed to decrease throughout aging, although
this could not be described quantitatively due to the
small precipitate size preventing SEM analysis, and
unknown specimen thickness preventing volume
fraction measurements by TEM. This dissolution of tertiary
b Fig. 1— STEM images of extraction replicas of alloy in the
as-homogenized state (a, b) and after 1 (c, d), 2 (e, f), 16 (g, h), 100 (i, j),
200 (k, l), and 1000 (m) h of aging at 1033 K (760 °C). Secondary γ′
precipitates are shown on the left-hand side (a, c, e, g, i, k, m) and
tertiary γ′ precipitates on the right-hand side (b, d, f, h, j, l). No
tertiary precipitates were visible after 1000 h of aging.
γ′ is driven by Oswald ripening, enabling the growth of
the larger secondary γ′ precipitates.
The APT reconstructions showed tertiary γ′
distributions consistent with those obtained by STEM imaging
of the carbon replicas, although statistical analysis of
the PSDs was limited by the low number of precipitates
in each APT needle. An example of one of the APT
reconstructions, from the sample aged for 16 hours, is
shown in Figure 5. In this figure, each green spot
represents a Cr ion and each blue spot represents an Al
ion. Isosurfaces corresponding to Cr concentrations of
15 at. pct. have been included to delineate the
matrix/precipitate boundaries. A large range of tertiary
γ′ sizes can be seen, along with a region bounding a
secondary γ′ precipitate in which no tertiary γ′
precipitates are present.
The compositional variation of each phase as a
function of aging time is shown in Figure 6. The
STEM–EDX data are presented as open markers, with
error bars corresponding to the standard deviation of
the values obtained. APT data are presented as solid
lines between solid markers and the ThermoCalc
predictions of the equilibrium compositions are given as
Measurements of the matrix composition, Figure 6(a),
indicated that it contained high concentrations of Cr
and Mo, as well as low levels of Al and Ti, relative to the
nominal alloy composition. These results were
consistent with the established elemental partitioning behavior
between the phases of commercial Ni-based
] There is good agreement between the
experimentally determined STEM–EDX and APT data for all
elements, with differences typically lying within the
range of experimental uncertainty. With increasing
aging time, the concentration of each element was
observed to be approximately constant and generally
corresponded well with the equilibrium concentrations
predicted by ThermoCalc. Of the elements present, only
the Cr concentration was consistently different to the
ThermoCalc predictions, being underpredicted by
∼2 at. pct.
The variation in elemental concentration with aging
time in the secondary γ′ is shown in Figure 6(b). As
expected, the Al and Ti concentrations are
comparatively high compared to the nominal alloy composition,
while the Cr and Mo levels are comparatively low. As
with the γ matrix composition, there is generally good
agreement between the elemental concentrations
measured by STEM–EDX and those obtained with APT.
However, a notable disparity exists between the Al
concentrations measured with the two techniques. This
is due to the fact that the Kα emissions from Al have
very low energy and are therefore easily absorbed by the
sample without being detected during STEM–EDX,
resulting in an underestimation of the Al
] As a consequence, where the experimental data
for Al differed in this way, the APT data were deemed to
be more reliable.
Progressive variations in elemental concentrations
with aging time in the secondary γ′ phase can be
discerned from the data presented in Figure 6(b). The Al
APT data suggest that the concentration of this element
increases from 11.3 at. pct. after 1 hour of aging to 12.5
at. pct. after 1000 hours. These values bound the
equilibrium concentration of ∼12 at. pct. predicted for
the γ′ phase by ThermoCalc. The Cr and Ti show
evidence of a decrease in concentration with increasing
aging time, although these changes are small and were
not significantly beyond the experimental uncertainty.
While the measured Cr concentration trended towards
that predicted by ThermoCalc, the Ti concentration
differed markedly, being approximately 1 to 2 at. pct.
below the predicted values. Notably, the concentration
Fig. 5— APT reconstruction of the alloy after 16 h of aging at 1033 K
(760 °C). Green spots represent the detection of one Cr ion, while blue
spots represent one Al ion. γ′ precipitates have been highlighted by the
addition of an isosurface at 15 at. pct. Cr. One secondary and many
tertiary precipitates are visible (Color figure online).
of Mo in the secondary γ′ increased from 0.7 at. pct.
after aging for 1 hour to 1.5 at. pct. after 1000 hours.
This change in Mo concentration was unexpected as it
tended away from the predicted concentration of ∼0.1
Analyses of the elemental concentrations in the
tertiary γ′, Figure 6(c), also identified unexpectedly high
levels of Mo, which were in excess of 2 at. pct. and
varied from 2.9 at. pct. after 1 hour of aging to 2.1
at. pct. after 100 hours of aging from the APT data.
Both the Al and Ti contents increased slightly on aging
to 100 hours, although the Ti content was markedly
lower than that of the secondary γ′. Significant
differences were observed between the Cr concentrations
measured by APT and STEM–EDX. The APT data
were consistent in the range of 4 to 6 at. pct., while the
STEM–EDX data varied from 5 to 14 at. pct. While it
may be possible that these discrepancies arose as a result
of statistical variations between the precipitates
sampled, the effect of retained γ within the carbon replica
samples cannot be entirely dismissed and would account
for the anomalously high Cr levels in the STEM–EDX
measurements at 1, 2, and 200 hours. With the exception
of Al, the equilibrium elemental concentrations
predicted by ThermoCalc do not agree well with the
experimentally measured values, nor do they appear to
be trending toward the predicted values with prolonged
time at temperature.
Figures 7(a) and (b) show the proxigrams of the
elemental concentrations across the secondary and
tertiary γ/γ′ interfaces measured by APT of the sample
aged for 16 hours. In these figures, all distances are
quoted with respect to the 15 at. pct. Cr isosurface.
Proxigrams for the other aging times are available in the
Supplementary Information. The elemental
concentraFig. 7— Proximity histogram of elemental concentrations across the
secondary γ′/γ interface (a) and the tertiary γ′/γ interface (b) for the
alloy aged for 16 h. (c) The variation of the secondary γ′/γ interface
width as a function of aging time (Color figure online).
tion profiles all showed sigmoidal variations. As with
previous reports, an accumulation of nickel was
observed ∼2 nm into the γ matrix.[
] The quantitative
analysis of the secondary γ′ interface width for each
element was performed in WaveMetrics Igor Pro and
the results are presented in Figure 7(c). Values for the
interface width for each element were calculated as the
distance between 1 pct above and 1 pct below the
minimum and maximum compositions, respectively.
These data showed a progressive decrease in the width
of the interface for all elements with increasing aging
time, from which it can be seen that Mo has the largest
interface width and Ni has the smallest.
The lattice parameters of the γ and γ′ phases and the
associated lattice misfit obtained by neutron diffraction
are presented in Figure 8. These results indicate that the
lattice parameters and lattice misfit are approximately
constant throughout aging, within the experimental
uncertainty. This is consistent with the nearly constant
APT compositions of the γ and γ′ phases, Figure 6. The
average measured lattice misfit was determined to be
−0.050 ± 0.008 pct. In comparison, the lattice misfit
predicted using ThermoCalc was 0.43 pct. The
inconsistency of this prediction with the experimental results
cannot be attributed to the effect of lattice constraint
between the two phases, since the signs of the predicted
and measured lattice misfits differ. Similarly, the
discrepancy cannot be ascribed to differences between the
measured and predicted compositions, as a larger,
positive lattice misfit is predicted when using the
experimentally determined compositions within
A. Precipitate Morphology
During the early stages of aging at 1033 K (760 °C),
the secondary γ′ size remained approximately constant,
while the volume fraction increased by ∼7 pct., Figures 3
and 4. This suggests that this distribution is in a growth
phase, leading to an increased number density of
precipitates. The apparent absence of coarsening during
this period differs from previous studies of precipitate
evolution in other superalloys[
] and could be a
consequence of the low lattice misfit, and associated
strain energy, between the γ and γ′ phases providing a
minimal driving force for this process. A constant lattice
misfit has previously been observed in a similar alloy on
aging up to 12 hours and linked to a constant secondary
γ′ precipitate size.[
After 100 hours of aging, the secondary γ′ showed
evidence of morphological changes, with cuboidal and
octodendritic precipitates being observed in the
microstructure. In alloys with moderate to high lattice
misfits, morphological changes of this type are
commonly associated with the magnitude of the lattice
misfit, with a progression from spheres to cuboids and
finally to octodendrites, in order to minimize the elastic
strain energy of the matrix as the precipitates
] However, in this alloy, the lattice misfit is low,
and therefore the precipitates would be expected to
remain spheroidal throughout aging.
The splitting of octodendritic precipitates into smaller
spherical or cuboidal precipitates is also driven by a
reduction in the overall energy of the system and occurs
when the overall elastic strain energy overcomes the
Fig. 8— Variation of the lattice parameter of each phase, and the
lattice misfit of the alloy as a function of aging time at 1033 K
(760 °C). The average lattice misfit during aging is shown as a
dotted line (Color figure online).
increased interfacial energy of the split precipitates.[
Although splitting is commonly associated with high
] it has been observed previously in low
misfit alloys such as RR1000 and UCO1.[
morphological instabilities that can lead to precipitate
] are known to occur in alloys of varying
misfit. It must be noted, however, that the mechanism
leading to these instabilities differs depending on the
magnitude of the lattice misfit. Alloys with low lattice
misfits may form precipitates with such morphologies as
a result of diffusion-induced interface instabilities.[
In contrast, in alloys with larger misfits, the process is
governed almost entirely by the strain energy.
Evidence of this process was observed in the
microstructure after aging for 100 hours, Figure 1(i).
Some regions of this sample contained large,
octodendritic precipitates indicative of the precursor to splitting,
while other regions contained only smaller, spherical/
cuboidal secondary γ′ precipitates, characteristic of
splitting having occurred. Splitting does not occur
instantaneously across the whole sample, and hence
the 100-hour aged sample contained both larger and
smaller secondary γ′ precipitates, resulting in a wider
particle size distribution. Additionally, the regions with
smaller, spherical/cuboidal precipitates exhibited locally
lower volume fraction than those containing larger
octodendritic precipitates. This accounts for the large
uncertainty associated with the volume fraction data
from this condition, Figure 4.
In addition, after the onset of splitting, there appeared
to be a decrease in the amount of tertiary γ′ visible in the
STEM images of the samples. This is consistent with
] However, quantification of the
tertiary γ′ volume fraction could not be reliably
performed from the analysis of the STEM images acquired
due to difficulties in accurately determining the local foil
thickness. The splitting phenomenon also serves to
explain the reduced size of the secondary γ′ after 1000
hours, with these precipitates having completed the
splitting process. Importantly, the morphological
changes observed are not consistent with those that
would be expected from instabilities governed by fast
growth in selected directions due to supersaturation of
the matrix phase.[
The tertiary γ′ precipitates were observed to increase
in size during aging, Figure 3, approximately following a
t1/3 relationship, as would be expected from classical
Lifshitz–Slyozov–Wagner (LSW) theory.[
However, using this relationship and extrapolating the
coarsening of the tertiary γ′ to 1000 hours, the predicted
size would remain below that of the secondary γ′
precipitates. This confirms that the absence of tertiary
γ′ after 1000 hours of aging is due to resolution of the
precipitates, as opposed to coarsening to a size sufficient
to merge with that of the secondary γ′.
B. Precipitate Composition
In the as-homogenized condition and in the majority
of the aged samples, the model alloy exhibited a bimodal
size distribution of secondary and tertiary precipitates.
Both the secondary and tertiary γ′ form on cooling from
the homogenization temperature, with the secondary γ′
forming before the tertiary γ′, at higher
] The higher elemental mobility during the
formation of the secondary γ′ produces larger
precipitates and facilitates the movement of Mo and Cr out of
the γ′ and into the matrix phase in which they
preferentially reside. In contrast, as a result of the lower
temperature and therefore reduced diffusion during the
formation of the tertiary γ′, this distribution is generally
enriched in those elements that partition to the γ phase
during secondary formation, resulting in a composition
further from the equilibrium predicted value. The Mo
content of the γ′ precipitate phase has been shown to
decrease with increasing aging temperature, in
agreement with this principle.[
During aging of the alloy, relatively modest changes
in the composition of the phases present were observed,
Figure 6. Most notably, the tertiary γ′ phase contained
markedly higher concentrations of Mo and Cr and a
lower concentration of Ti than the secondary γ′
precipitates. While the Mo and Al concentrations in the two
precipitate size distributions appear to be converging
with increasing aging time (as seen in Reference 43), it is
not clear that this is true of the Cr and Ti
concentrations. The elevated Mo concentrations observed in the γ′
precipitates are consistent with data reported by Jou
] on IN100. Similarly, a significant Mo
concentration in the secondary γ′ phase has also been reported
by Loomis et al.[
] in an investigation of similar model
alloys to those studied here. However, a previous study
comparing experimental compositions from APT to
equilibrium predictions by ThermoCalc did not find the
Mo content in the γ′ to be as high as the present
It has been shown previously that the composition of
γ′ precipitates is dependent on size within the range of
5 nm to 3 μm.[
] In the present study, no link was found
between the size of the γ′ precipitates within each size
distribution and the composition.[
] However, the range
of precipitate sizes within each distribution in the
current study is small and, therefore, a wider range of
precipitate sizes may have revealed compositional
The concentration profile widths in Figure 7(c)
decreased continuously during aging for all elements.
This is consistent with the studies of Ni-Al,[
] and Ni-Al-Cr-Ta[
] alloys, which showed rapid
decreases in the interface widths in the early stages of
aging and lower rates of decrease with prolonged time at
temperature, as the precipitates entered the coarsening
regime. However, it is noted that the opposite behavior
of increasing interfacial width has been reported during
the aging of secondary γ′ precipitates in the commercial
superalloy Rene 88DT.[
] This suggests that a
number of factors may influence interfacial width. For
example, it has been shown that a sharper interface
results from slower cooling rates in AM1. The relative
magnitudes of the interfacial width differ for each
element, with the faster diffusing species, Al and Ti,
exhibiting narrower interfaces than those that diffuse
more slowly. This is particularly notable for Mo, which
displays the widest interface despite having the smallest
difference in concentration between the γ and γ′ phases.
Theoretical descriptions exist that account for the finite
width of the γ/γ′ interface, as observed in this study.
These include the original work by Cahn and Hilliard[
that utilized gradient energy coefficients to account for
the free energy of regions of non-constant composition,
and the more recent trans-interface chemical diffusion
model by Ardell[
] where the kinetics are controlled
by diffusion through, rather than to, the γ/γ′ interface.
Composition has previously been linked to precipitate
morphology in a study on the low misfit alloy RR1000,
where it was shown that restricted elemental diffusion
resulted in cyclic coarsening and splitting of the
] It was shown that the Al concentration of
the γ′ phase generally increased on aging at 1073 K
(800 °C), reaching a maximum immediately following
splitting. In contrast, the concentration of Ti remained
approximately constant during aging but exhibited a
local minimum after splitting. The overall trends in the
composition of Al and Ti in the alloy here studied are
not inconsistent with these results, although sampling at
more regular intervals would be required to determine
the local changes associated with splitting.
C. Lattice Misfit
The difference between the secondary and tertiary γ′
compositions would be expected to result in differing
lattice parameters for these distributions and,
consequently differing lattice misfits. However, the broad
intrinsic peak width and comparatively low volume
fraction of the tertiary γ′ phase prohibited separation of
the contributions of this distribution to the superlattice
reflections in the neutron diffraction data, which were
therefore dominated by the secondary γ′.[
D. Consequences for Mechanical Properties
The differences between the measured compositions
of the γ′ precipitates and those predicted by ThermoCalc
have implications for the estimation of alloy strength
using empirical relations.[
] Specifically, the high
Mo and low Ti concentrations will affect multiple
mechanisms of alloy strengthening concurrently,
including solid solution, coherency, and order strengthening.
Consequently, the contribution from each strengthening
mechanism will vary between the secondary and tertiary
γ′ populations as a result of their different compositions.
Therefore, relationships describing the properties of
superalloys may require modification to accommodate
the compositional differences between the γ′
populations, to further improve their fidelity.
The microstructure and phase composition of a model
polycrystalline Ni-based superalloy have been
characterized using STEM–EDX and APT, following aging at
1033 K (760 °C) for durations of up to 1000 hours.
Microstructural analysis revealed that the size of the
secondary γ′ precipitates remained approximately
constant during the early stages of aging, but exhibited the
morphological instabilities associated with precipitate
splitting following longer duration exposures. In
contrast, when the tertiary γ′ precipitates were present in the
microstructure, they were observed to coarsen in line
with LSW theory. The experimentally determined phase
compositions were compared to equilibrium predictions
made using ThermoCalc. The composition of the γ
phase was seen to remain approximately constant
during aging, with elemental concentrations similar to
the equilibrium phase composition predicted by
ThermoCalc. Compositional analysis of the secondary and
tertiary γ′ precipitates revealed small changes in
elemental concentration with increasing aging time. Notably,
higher Mo and lower Ti concentrations were observed
than those predicted using ThermoCalc. These results
are believed to have implications for models describing
alloy properties, which currently employ
thermodynamic equilibrium predictions of phase composition.
The authors wish to acknowledge Mrs. S. Rhodes,
Dr. H.T. Pang, Dr. D.M. Collins, and Dr. O.M.D.M.
Messe´ for their assistance with the experiments
performed. Funding was provided by the
EPSRC/RollsRoyce Strategic Partnership under EP/M005607/1 and
EP/H022309/1. The Oxford Atom Probe facility was
funded by the EPSRC under EP/M022803/1. Neutron
diffraction beam time was supported through the
Canadian Neutron Beam Centre under Experiment
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