Characterization of Phase Chemistry and Partitioning in a Family of High-Strength Nickel-Based Superalloys
Characterization of Phase Chemistry and Partitioning in a Family of High-Strength Nickel-Based Superalloys
M.T. LAPINGTON mark.lapington@materials 0
D.J. CRUDDEN 0
R.C. REED 0
M.P. MOODY 0
P.A.J. BAGOT 0
0 M.T. LAPINGTON, M.P. MOODY and P.A.J. BAGOT are with the Department of Materials, University of Oxford , Parks Road, Oxford OX1 3PH , UK. Contact
A family of novel polycrystalline Ni-based superalloys with varying Ti:Nb ratios has been created using computational alloy design techniques, and subsequently characterized using atom probe tomography and electron microscopy. Phase chemistry, elemental partitioning, and c¢ character have been analyzed and compared with thermodynamic predictions created using Thermo-Calc. Phase compositions and c¢ volume fraction were found to compare favorably with the thermodynamically predicted values, while predicted partitioning behavior for Ti, Nb, Cr, and Co tended to overestimate c¢ preference over the c matrix, often with opposing trends vs Nb concentration.
NICKEL-BASED superalloys are widely used in the
high-pressure section of gas turbine engines due to their
exceptional oxidation resistance and strength at
temperatures up to and exceeding 1000 C.[
and environmental drivers seek to increase turbine
operating temperatures, requiring further alloy
optimization in order to satisfy the increasing demands on
strength and oxidation resistance. Complex trade-offs
have to be considered in order to retain these properties
alongside fatigue resistance, creep resistance, and phase
stability during the optimization process. Much of the
high-temperature strength of these alloys come from the
precipitation of fully coherent ordered FCC c¢ phase
] These precipitates retain strength at high
temperatures via their unique ability to resist
deformation by promoting cross slip of partial dislocations onto
] This deformation mechanism is heavily
influenced by c¢ volume fraction and chemistry, as c¢
forming elements such as Ti, Ta, and Nb have been
shown to increase strength by raising the stacking fault
energy at anti-phase boundaries (APB) found in the
] However, there are limits
regarding the addition of c¢ forming elements, as high
c¢ volume fraction has been shown to have detrimental
effects on processing and fatigue resistance,[
addition of Nb and Ti can adversely affect c¢ phase
stability due to their potential to form deleterious
secondary phases such as d (Ni3Nb)[
] or g (Ni3Ti).[
Ti has also been previously shown to reduce oxidation
resistance in model Ni-Cr-Ti polycrystalline alloys.[
Optimization of Ti and Nb levels is therefore
necessary in order to find an appropriate balance between
high-temperature strength, oxidation resistance, and
phase stability. Recent advances in computational
optimization techniques have utilized thermodynamic
modeling to predict phase chemistries and physical
properties based on composition alone, resulting in
major cost and time savings. The Alloys-by-Design
(ABD) process developed by Reed et al.[
] utilizes a set
of merit-indices to sort through these compositions
looking to balance the trade-offs mentioned above. This
ABD process has been modified for application to
polycrystalline superalloys by Crudden et al.,[
has been used to create a family of superalloys to
investigate composition-property relationships, with a
focus on the effect of Ti and Nb on microstructure and
phase chemistry as they relate to oxidation resistance
and high-temperature strength.
The small size of c¢ precipitates and need for accurate
phase chemistry determination requires a quantitative
analysis technique with high spatial resolution and
elemental sensitivity. Atom probe tomography (APT)
combines both of these with powerful 3D analysis
] and has been successfully used in the
] to investigate phase chemistries[
partitioning ratios in nickel-based superalloys. The
high spatial resolution and 3D analysis also allows
detailed investigation of nanoscale tertiary c¢
precipitates, as well as c-c¢ interfaces.[
To study the effect of varying Ti:Nb ratio, a family of
superalloys has been created which are near identical in
composition, except for the substitution of Nb for Ti on
a 1:1 basis. This similarity of composition between the
alloys allows direct comparisons of microstructure and
chemistry to be made. The target alloy compositions[
given in Table I result from using the ABD method to
predict compositions with high strength at 800 C
combined with good oxidation resistance. The material
that was examined in this investigation was extracted
from pancake forgings, which were isothermally forged
from hot isostatic pressed powder. The as-manufactured
alloy compositions are also provided in Table I, which
were measured using a combination of X-ray
fluorescence (XRF) spectroscopy, the combustion infra-red
method and inductively coupled plasma—optical
emission spectroscopy (ICP-OES). A two-stage super-solvus
heat treatment was carried out, including c¢
solutionizing at 1170 C for 2 hours followed by aging at 850 C
for 4 hours.
Metallographic samples for SEM analysis were
sectioned using an Isomet 5000 diamond cutting wheel,
then ground and polished with successively finer grades
of SiC grit paper and diamond suspension. A mirror
finish was achieved using 0.06 lm COL-K colloidal
silica, and no subsequent etching was performed. SEM
images have been captured using a Zeiss Nvision 40
FIB-SEM with 5 kV beam current to a maximum
magnification of 20,000 times, using the inlens
secondary electron detector to enhance the visual contrast
between c and c¢ phases. ImageJ software[
] has been
used to process the SEM images and calculate the
cross-sectional area and size distribution of c¢
Sample blanks for atom probe tomography were cut
from sections of superalloy into 0.5 mm 9 0.5 mm 9
20 mm cuboidal rods. A two-stage electropolishing
process was used to create needle-shaped samples with
a tip diameter < 100 nm,[
] with 25 pct perchloric acid
solution (in acetic acid) at a voltage of 16V for first-stage
shaping, and 2 pct perchloric acid solution (in
butoxyethanol) at a voltage of 10 V for second stage
polishing. Specimens were analyzed in a Cameca LEAP 3000X
HR in laser-pulsing mode, at a pulse frequency of
200 MHz, laser pulse energy between 0.4 and 0.5 nJ,
and a stage temperature of 50 K. Detection rates (the
average percentage of pulses resulting in ion detection)
were targeted around 2 pct in order to maximize yield.
All atom probe data have been reconstructed and
analyzed using Cameca IVAS 3.6.12 software. Peaks
observed in the mass spectra were ranged using a
background-to-background ranging technique, with
each individual range spectra templated from a universal
ABD-superalloy range file to ensure consistency. The
freely available APTtools software suite[
routinely referenced to ensure that elemental peak
signatures identified during ranging are accurate and
complete. Peak overlaps consisting of multiple elemental
species have been deconvolved using standard IVAS
decomposition parameters, which helps to maintain
quantitative composition accuracy at the potential cost
of spatial information. Phase compositions were
measured using 10 nm 9 10 nm 9 10 nm volumes located
far from phase boundaries. The average composition of
two volumes was used to calculate matrix composition
(16 volumes totalling 5.5 9 105 ions), and c¢ precipitate
compositions were given as a weighted average of all
precipitates in the sample (32 volumes totalling
5.9 9 105 ions).
Thermodynamic modeling was carried out in
Thermo-Calc 2015b using the TTNi8 (Thermo-Tech) and
TCNi8 (Thermo-Calc) databases. Phase composition
calculations were based on the as-manufactured
compositions, and performed at the final alloy aging
temperature of 850 C. The superalloys have been
assumed to be in a 3-phase system consisting of c, c¢
and MC carbides, which were the only phases observed
at room temperature during X-ray synchrotron
] with all other phases suspended.
A. Scanning Electron Microscopy
Several SEM micrographs were taken at two
magnifications from each of the studied alloys as shown in
Figure 1. The micrographs taken at 500 times
magnification revealed average grain sizes between 25 and
40 lm, which was within the target size range selected
for a balance between yield strength, creep strength, and
crack propagation resistance. Serrated grain boundaries
and twinning were also observed in each alloy. At a
magnification of 20,000 times, intragranular secondary
and tertiary c¢ precipitates were clearly visible, with the
secondary precipitates occurring as a mix of spheroidal
and cuboidal/octodendritic morphologies, and the
smaller tertiary precipitates appearing to be spheroidal
in all alloys. In contrast, the precipitates in the high-Nb
alloy ABD-6 (Ti2.8Nb1.2) showed a unimodal
distribution with a spheroidal morphology. No primary c¢ was
observed at the grain boundaries, so all precipitate
analysis was performed within grains.
Further analysis was carried out using the ImageJ
software package to determine the c¢ cross-sectional area
fraction and size distribution. Micrographs from several
areas per sample were converted into 8-bit grayscale
images, from which brightness histograms were created
showing two peaks representing the dark c matrix and
lighter c¢ precipitates. The c-c¢ interface was defined by
selecting a threshold brightness value located at the local
brightness minima between the two peaks to ensure
consistency across the analyses. The ImageJ particle
analysis tool was employed to measure the area of each
detected precipitate in the thresholded images. The effect
of noise was reduced by discounting precipitates with an
area < 300 nm2 (equivalent diameter 19.5 nm), which
represents any precipitate covering < 9 pixels of the
SEM micrograph at 20,000 times magnification.
Individual precipitate areas were summed to find the
total c¢ cross-sectional area fraction for each
micrograph. This can be considered to equal the precipitate
volume fraction as long as Delesse-Rosiwal conditions
are met, which requires the surface to lie on a two
] For this reason, the samples were
kept unetched, and a low electron beam current was
used to minimize interaction depth. Three micrographs
were taken for each alloy, resulting in mean average c¢
volume fractions of 48, 48, 43, and 47 pct for Alloys
ABD-2, 4, 5, and 6, respectively (as summarized in
Table III). The individual precipitate areas were also
used to determine the precipitate size distributions by
converting particle cross-sectional area (A) inptoffiffiffiffiffiaffiffiffinffi
equivalent diameter (d) using the formula d ¼ 2 A=p.
Figure 2 shows the total precipitate size distributions
summed over all micrographs of each alloy. ABD-2, 4,
and 5 alloys (with compositions Ti4.1Nb0, Ti3.6Nb0.4,
and Ti3.2Nb0.8, respectively) all show a large peak at
40-nm equivalent diameter, which corresponds to the
smallest tertiary c¢ precipitates not discounted as noise.
Secondary c¢ precipitates appear to have a bimodal
distribution with precipitates in ABD-2 and 4
averaging 140 or 260 nm diameter, while ABD-5 shows
unimodal secondaries of average 180 nm. The high-Nb
alloy ABD-6 (Ti2.8Nb1.2) is the outlier here, with very
few tertiary precipitates and a unimodal secondary
distribution peak centered on 160 nm.
B. Atom Probe Tomography
A total of eight successful atom probe experiments
were run, comprising two specimens per alloy, with
reconstructions varying in size from 26 to 105 million
ions. The microstructure present in each of the alloys is
shown in Figure 3, which shows 10-nm-thick sections
clipped through the center of atom maps. Only Ti (blue)
and Cr (pink) ions are shown as these elements partition
strongly to c¢ precipitates and c matrix phases
respectively, allowing easy visual identification. The phase
preference of each major constituent element is
demonstrated in Figure 4, which consists of a 10-nm-thick slice
clipped from a reconstruction of the ABD-4 alloy split
into four pairs of elements. All elements were found to
partition as expected, with Ni, Al, Ta, Ti, and Nb all
partitioning to c¢ precipitates while Cr and Co partition
to the c matrix. Several generations of intergranular c¢
precipitate were observed in reconstructions, occurring
with a bimodal distribution of larger secondary c¢II
(> 100 nm) and smaller spheroidal tertiary c¢III (20 to
The overall composition for each sample was
measured using the total ranged ion counts, from which a
mean average composition was obtained for each alloy
as shown in Table II. The boundary between c and c¢
phases was defined with a 22.5 at. pct chromium
isoconcentration surface, and the composition of each
phase was found using a 10 nm 9 10 nm 9 10 nm
region of interest far from this interface. The phase
compositions presented in Table II are a weighted
average of all occurrences of each phase across datasets
from a particular alloy. Both secondary and tertiary c¢
precipitates have been included in the average c¢ phase
composition as there was no observed correlation
between precipitate size and composition.
The c-c¢ interface was also investigated for potential
signs of segregation. Figure 5 shows two concentration
profiles calculated as a function of the distance to a
22.5 at. pct Cr isosurface, also known as a proximity
]). These were taken over a
10-nm region crossing the interface in the ABD-5 alloy,
separated into major (> 5 at. pct nominal) and minor
(< 5 at. pct nominal) constituent elements. There is a
2-nm-wide transition region between c and c¢ phase
compositions for most elements, with some notable
exceptions. Ti and Ta do not reach equilibrium c¢
concentration for several nm into the precipitates, while
the Cr concentration entering the c matrix seems to
plateau before rising again 2 nm from the interface.
Additionally, W enrichment can be seen directly at the
c-c¢ interface, present at double the nominal W
concentration (Wnom = 0.8 at. pct). These features have been
observed in multiple proxigrams across c-c¢ interfaces
for each of the studied alloys.
Determination of phase chemistries allows an
estimate of the c¢ volume fraction /c¢ to be made using a
derivation of the lever rule[
Cn ¼ /c0 Cc0 þ ð1
where Cn, Cc, and Cc¢ are the concentrations of an
element present in the nominal alloy composition, and
measured in the c and c¢ phases, respectively. The c¢
volume fraction can be estimated by plotting Cn Cc vs
Cc¢ Cc for each element and measuring the gradient of
the best fit line, as shown in Figure 6. c¢ volume
fractions obtained using this technique is shown in
Table III, combined with volume fractions observed
during SEM and predictions made with Thermo-Calc
using both TTNi8 and TCNi8 thermodynamic
databases. The observed volume fraction values tend to lie
very close to, or even between the two values predicted
by Thermo-Calc using both databases.
Partitioning characteristics for each element are
determined from Figure 6, with c¢ forming elements
located in the top right quadrant and c formers in the
bottom left. Elements centered about the origin have
been magnified in inset (b). Most elements fall directly
on the line of best fit, as the total composition is simply
the sum of the c and c¢ phase contributions. Exceptions
to this include C, B, and Zr which are all known to
segregate to grain boundaries, and W which has been
observed at elevated levels at c-c¢ interfaces. The phase
partitioning preference of a particular element can be
quantified using the partitioning coefficient Kc¢/c, defined
as the concentration of that element in the c¢ phase
divided by the concentration in c phase. Consequently,
high Kc¢/c values indicate strong partitioning to c¢ and
near-zero Kc¢/c values partition strongly to c. Figure 7
displays the partitioning coefficients of Ti, Nb, Cr and
Co across the four alloys, with Kc¢/c values measured via
APT (shown in blue) compared with values predicted by
Thermo-Calc (red using TCNi8 database, green using
TTNi8). The observed Kc¢/c values were found to be
consistently lower than values predicted using the
TTNi8 database, resulting in a lower concentration of
these elements present in the c¢ phase. This is made up
for by an increase in the observed c¢ partitioning of Ni.
Predictions made using the TCNi8 database were much
closer to observed values, but mostly showed opposing
trends for Kc¢/c vs Nb concentration. In addition,
predicted Kc¢/c values for Ti, Nb, and Co in alloy
ABD-5 show a slight reduction in partitioning strength
compared to the linear trendlines, which may result
from the dip in Al concentration observed in the
as-manufactured compositions in Table I.
A. c¢ Microstructure
Precipitation of c¢ phase is one of the most crucial
strengthening mechanisms in these particular Ni-based
superalloys, for which the volume fraction and size
distribution are important factors. The alloys studied in
this report were designed to have a c¢ volume fraction of
around 50 pct, which has been verified using both APT
and SEM analysis as recorded in Table III.
Interestingly, both experimental techniques and the two
thermodynamic databases predict a slightly lower c¢
volume fraction for alloy ABD-5 compared to the other
alloys. It is tempting to attribute this lower c¢ volume
fraction to the slight Al deficit seen in the
as-manufactured compositions in Table I (7.6 at. pct measured vs
8.0 at. pct target), which could lead to a reduction in the
nominally Ni3Al c¢ phase. However, the difference in c¢
volume fraction is an order of magnitude larger than the
Al compositional deficit, so no solid conclusions can be
drawn at this time.
Most of the ABD alloys express this c¢ volume
fraction as a mixture of small tertiary precipitates in
addition to a bimodal distribution of secondaries, which
is backed up by research on the same ABD-series alloys
carried out by Hisazawa et al.[
] In the as-received state,
the alloys were reported to have a very small fraction of
primary c¢ due to the super-solvus heat treatment
regime, plus secondaries in the ~ 200nm size range and
tertiaries < 50 nm, which almost perfectly matches the
SEM observations and size distributions shown in
Figures 1 and 2. The only discrepancy involves the
high-Nb ABD-6 alloy, which was observed in this study
to have a unimodal size distribution with little
differentiation between secondary and tertiary precipitates. It
seems unlikely that this can be directly linked to the
increased Ti/Nb ratio in the ABD-6 alloy, as Hisazawa
et al. saw no correlation between Ti/Nb ratio and size
distributions, likely due to the similar diffusion rates of
Ti and Nb leading to similar nucleation and growth
kinetics. The unimodal distribution combined with the
change to a more spherical morphology indicates that
the ABD-6 sample has undergone higher rates of growth
than nucleation during heat treatment, possibly due to
its position within the forging.
B. Interface Chemistry
One of the unexpected results from this experiment was
the discovery of W enrichment at c-c¢ interfaces. The first
indication came from the Lever rule plots used to
calculate c¢ volume fraction (an example of which is
shown in Figure 6), where the plotted positions of W, C,
B, and Zr data points did not lie directly on the best fit line.
C, B, and Zr are known to preferentially segregate to grain
boundaries, so investigation of interface regions was
carried out. None of these grain boundary segregants
were found, but a thin shell of W-enriched material
surrounding c¢ precipitates was revealed. This behavior
has also been seen by Amouyal et al. in ME-9 alloy[
a single crystal directionally solidified superalloy.[
] It is
unknown if this behavior is due to preferential
segregation, or rejection of W during c¢ formation. The shape of
the W concentration profile over the c-c¢ interface is
similar to the ‘bow-wave’ profile seen by Warren et al.[
for Re in single-crystal RR3000, which was attributed to
diffusion-limited rejection of solute. This could also apply
to W, which has been shown by Karunaratne et al.[
have a similar, but slightly larger interdiffusion coefficient
to Re through c phase.
The Ti and Ta concentration profiles were also found
to be non-uniform over the c-c¢ interface, with both
elements continuously increasing in concentration up to
5 nm within the c¢ precipitates. A similar effect was
observed for Ti, Ta, and Hf in the RR1000 alloy as seen
by Bagot et al.,[
] where it was attributed to the large
lattice misfits for these elements as estimated using
] Cr concentration profiles were also
found to be non-uniform heading into the c matrix
phase, showing a trend of rising to a plateau
concentration of roughly 30 at. pct, before rising another 2 to
4 at. pct 2 nm from the interface. This trend was
repeated across 29 of the 32 analyzed c-c¢ interfaces.
The origin of this slight non-uniformity requires further
detailed study, as it is possible that it may be a field
evaporation artefact caused by moving through the
C. Alloy Composition/Phase Chemistry
In addition to volume fraction, control of alloy
composition is critical to ensuring a good balance of
properties within both the matrix and precipitates. The
example of reduced Al content in the as-manufactured
ABD-5 composition mentioned above has slight
knock-on effects for both precipitate volume fraction
(Table III), and the partitioning ratios of other
c¢-forming elements (Figure 7), showing how slight changes can
influence these chemically complex systems. One of the
reasons for creating these alloys is to assess whether Ti
can be swapped out for Nb without causing major
changes to the alloy properties. Both Ti and Nb
promote high-temperature strength by increasing APB
energy in the c¢ phase,[
] but the effect is lessened if they
partition to the matrix, which highlights the importance
of their partitioning characteristics. Cr and Co fulfil a
similar role in the matrix by providing solid solution
strengthening. The partitioning graphs shown in
Figure 7 for these four elements show that the addition
of Nb does have a minor effect, with Ti partitioning less
strongly to c¢, while Nb, Cr, and Co migrate into the c¢
These results have then been compared with the
partitioning coefficients predicted by Thermo-Calc using
two different databases—TTNi8 (Thermo-Tech) and
TCNi8 (Thermo-Calc). The TTNi8 values tend to
overestimate the c¢ preference exhibited by Ti, Nb, Cr,
and Co. This mirrors similar findings relating to
thermodynamic predictions of Nb partitioning by
Antonov et al.[
] and Miller et al.,[
] who notes that
overestimation of Nb c¢ partitioning can lead to inflated
The Kc¢/c vs Nb pct trendlines also oppose the
observed trends for the elements Nb, Cr, and Co. The
TCNi8 predictions fall much closer to the observed
partitioning values, but still show opposing trends for
Ti, Cr, and Co. These discrepancies may be due to the
assumption of thermodynamic equilibrium which is
never reached due to material inhomogeneities or
unaccounted kinetics and diffusion.[
modeling is a powerful tool which gives us the ability to
perform billions of potential experiments in silico, but it
must be acknowledged that these are still just
predictions, and used accordingly. However, this study has
shown that careful, methodical use of APT can play a
key role in benchmarking the results from theoretical
predictions, as also recently demonstrated on in-service
] This combination of APT and modeling thus
offers a robust, general methodology which should be
considered in alloy design.
V. SUMMARY AND CONCLUSIONS
1. SEM micrograph analyses have shown that Alloys
ABD-2 (Ti4.1Nb0), ABD-4 (Ti3.6Nb0.4), and
ABD-5 (Ti3.2Nb0.8) form both secondary and
tertiary intragranular c¢ precipitates, with a bimodal
distribution of secondaries. The high-Nb alloy
ABD-6 (Ti2.8Nb1.2) was shown to have a unimodal
secondary c¢ distribution with spheroidal
morphology for this particular sample.
2. Intragranular precipitate volume fractions have
been estimated using cross-sectional areal
fractions found via ImageJ analysis of SEM
micrographs. The observed volume fractions match
values found via APT using lever rule analysis,
and show good correlation with Thermo-Calc
predictions made using both TTNi8 and TCNi8
3. Atom-Probe Tomography has been used to create
representative 3D maps of the microstructure and
quantitatively measure the compositions of c phase
matrix and microscale c¢ precipitates, which
compare favorably with both composition predictions
4. Both secondary and tertiary generations of c’
precipitates were found to have similar
compositions, with no apparent size-dependence. W was
observed to segregate to c-c¢ interfaces of multiple
c¢ precipitates across several datasets, possibly
due to diffusion-limited pileup during c¢ phase
5. All elements were found to partition to the phases
expected, although partitioning coefficients for Ti,
Nb, Cr, and Co were found to differ from predicted
values. Kc¢/c values predicted using the TTNi8
database tended to overestimate c¢ partitioning
behavior, whilst values predicted using the TCNi8
database were much closer but often showed
opposing Kc¢/c vs Ti:Nb trends.
The authors would like to acknowledge the financial
support of the UK Engineering and Physical Sciences
Research Council (EPSRC) and the Rolls-Royce
Strategic Partnership in Structural Metallic Systems
for Gas Turbine Applications EP/H022309/1 and EP/
H500375/1 and Rolls-Royce plc for the provision of
material. The Oxford Atom Probe facility has been
funded by the EPSRC Grant EP/D077664/1. The
research materials supporting this publication can be
accessed by contacting the Oxford Research Archive at
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