Materials Science and Technology
ISSN: 0267-0836 (Print) 1743-2847 (Online) Journal homepage: https://www.tandfonline.com/loi/ymst20
Formation of secondary reaction zone in
ruthenium bearing nickel based single crystal
superalloys with diffusion aluminide coatings
D. K. Das, B. Gleeson, K. S. Murphy, S. Ma & T. M. Pollock
To cite this article: D. K. Das, B. Gleeson, K. S. Murphy, S. Ma & T. M. Pollock (2009) Formation
of secondary reaction zone in ruthenium bearing nickel based single crystal superalloys
with diffusion aluminide coatings, Materials Science and Technology, 25:2, 300-308, DOI:
10.1179/174328408X382352
To link to this article: https://doi.org/10.1179/174328408X382352
Published online: 19 Jul 2013.
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Formation of secondary reaction zone in
ruthenium bearing nickel based single crystal
superalloys with diffusion aluminide coatings
D. K. Das1, B. Gleeson2, K. S. Murphy3, S. Ma1 and T. M. Pollock*1
Several Ru bearing experimental single crystal superalloys were coated with b-NiAl and Pt–Hf
modified c–c9 coatings and subjected to elevated temperature exposure. The microstructure,
oxidation resistance and the propensity for the formation of secondary reaction zones (SRZ) were
investigated for these coatings. The presence of significant amounts of Ru in the superalloys did
not prevent the cellular transformation leading to the formation of SRZ beneath the b-NiAl coating.
Cyclic oxidation exposure of the b-NiAl coated alloys at 1100uC led to a significant further growth
of the SRZ. In contrast, the c–c9 coating did not induce the undesirable cellular transformation,
even after prolonged high temperature exposure. The c–c9 coating also provided good oxidation
resistance to the superalloys. The driving forces for SRZ formation in case of the b-NiAl coating as
compared to the c–c9 coating are discussed.
Keywords: Ru containing superalloy, Secondary reaction zone, b-NiAl coating, c–c9 coating, Platinum aluminide coating
Dedicated to the memory of Professor Malcolm McLean
Introduction
The requirement for enhanced elevated temperature
mechanical properties in Ni base single crystal superalloys has led to increasingly higher additions of
refractory elements such as Re, W and Mo. However,
such additions often lead to the precipitation of
detrimental refractory element rich topologically close
packed (TCP) phases during heat treatment and/or
service.1–3 In Re containing single crystal superalloys,
three types of Re rich TCP phases, namely tetragonal s,
rhombohedral m and orthorhombic P, have been
observed.4 These phases often form as precipitates
distributed throughout the single crystal substrate via
continuous precipitation.2,4 They can also form by
discontinuous precipitation, also known as cellular
transformation, where the parent two phase c–c9
structure of the superalloy is transformed to a three
phase cellular structure consisting of c, c9 and TCP
precipitates.3,5,6 The transformation, which is driven by
local reductions in chemical free energy and, in some
cases, strain energy,5,7,8 occurs through a combined
mechanism of boundary precipitation and interfacial
migration.7,8 High angle grain boundaries in polycrystalline alloys and defects such as freckles in single crystal
alloys serve as sites for heterogeneous nucleation of
1
Materials Science and Engineering, University of Michigan, Ann Arbor, MI
48109, USA
Mechanical Engineering and Materials Science, University of Pittsburgh,
Pittsburgh, PA 15261, USA
3
Howmet Research Corporation, 1500 South Warner Street, Whitehall, MI
49461, USA
2
*Corresponding author, email tresap@umich.edu
300
ß 2009 Institute of Materials, Minerals and Mining
Published by Maney on behalf of the Institute
Received 9 September 2008; accepted 11 September 2008
DOI 10.1179/174328408X382352
precipitates and also as high diffusivity mobile reaction
fronts.5 Such cellular transformation is also known to
occur beneath b-NiAl type diffusion aluminide coatings,
that are typically applied for high temperature oxidation
protection. The transformed region beneath the aluminide coating is often referred to as the secondary
reaction zone (SRZ). Both b-NiAl and Pt modified bNiAl coatings have been reported to induce SRZ
formation during aluminisation treatment as well as
during high temperature exposure in service.5,6 The
presence of SRZs in superalloys is highly undesirable
since it leads to degradation in creep properties, caused
by cracking along cell boundaries in the transformed
region.9
The relatively high Al and low Ni contents in b-NiAl
type coatings facilitate SRZ formation in Ni based
superalloys due to extensive interdiffusion10,11 Further,
differences in the coefficients of thermal expansion
between the b-NiAl coatings and the superalloy
substrate induce stresses during thermal cycling. In
order to overcome these drawbacks, coatings having a
c–c9 rather than b-NiAl structure are currently being
explored.12–15 The Pt modified c–c9 type diffusion
aluminide coatings have been reported to have good
chemical and mechanical compatibility with the substrate.12,13 Additionally, they also have been reported
to offer the desired oxidation resistance to the superalloy substrates.12 In the present study, a plain b-NiAl
coating and a Pt modified c–c9 type coating have been
studied on a set of experimental Ru bearing Ni based
single crystal superalloys. Two major aspects, namely
the coating microstructures and the propensity for
SRZ formation in presence of the coatings, have been
examined.
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Formation of secondary reaction zone in single crystal superalloys
Table 1 Composition of alloys used in present study, wt-% (at.-%)
Alloy
Ni
F-26
F-16
F-13
F-18
F-22
F-20
F-30
62.3
55.5
59.2
65.8
67.0
63.6
54.8
Al
(66.3)
(59.2)
(65.1)
(69.9)
(69.6)
(66.9)
(59.0)
6.0
6.2
5.6
6.0
6.0
6.0
6.0
Ru
(13.9)
(14.3)
(13.4)
(13.8)
(14.4)
(13.8)
(13.8)
5.7
9.7
14.1
5.7
5.7
5.7
5.7
Ta
(3.5)
(6.0)
(9.0)
(3.5)
(3.6)
(3.5)
(3.5)
8.0
6.5
6.3
8.0
8.0
8.1
8.0
Re
(2.8)
(2.2)
(2.2)
(2.8)
(2.9)
(2.8)
(2.7)
Experimental
The compositions of the seven experimental single
crystal alloys studied are provided in Table 1. The Ru
content of the alloys was relatively high at 5?7 wt-% or
more. These alloys were originally prepared for studying
the effect of Ru alloying on solidification characteristics
and mechanical behaviour.16,17 As evident from Table 1,
alloys F-26, F-16 and F-13, in that order, had increasing
contents of Ru. Similarly, alloys F-18, F-22 and F-20
had increasing concentrations of Cr. Alloy F-30 was also
studied because it exhibits the best creep properties
among the above alloys.17,18 It should be noted that
alloys F-13 and F-18 did not contain any Cr. Single
crystal alloy rods were solutioned at 1300uC for 8 h
followed by aging at 1100uC for 8 h, to develop the c–c9
structure of the substrate.16,17 Disc shaped samples of
about 15 mm diameter and 4 mm thickness were sliced
from the heat treated rods, polished with a 600 grit
emery paper, and subsequently cleaned thoroughly. The
b-NiAl coating was applied at Howmet Corp.
(Whitehall, MI, USA) using a low activity aluminisation
process at 1080uC for 9 h.19 The c–c9 coating was
applied using a method developed by Gleeson et al.13
which involves Pt electroplating, diffusion treatment and
subsequent codeposition of Al and Hf by a proprietary
pack cementation process.13
The coated samples were subjected to cyclic exposure
at 1100uC in air using an automated cyclic oxidation
furnace. Each one-hour cycle consisted of 10 min heatup, 45 min dwell (at 1100uC) and 5 min cooling by a fan.
The temperature of the samples reduced to about 100uC
after the cooling period. The cumulative dwell time at
1100uC was considered to be the exposure duration for
any oxidised sample. The cycling was carried out for a
maximum of 210 cycles, i.e. exposure duration of about
160 h.
The microstructures of the as coated and the oxidised
samples were examined using both an Olympus PME3
optical microscope and a Hitachi S3200N scanning
electron microscope (SEM) operating at 15–20 kV. An
etchant consisting of 33 vol.-% acetic acid, 33 vol.-%
nitric acid, 1 vol.-% hydrofluoric acid and 33 vol.-%
distilled water, was used to reveal the microstructural
details. Compositional information including X-ray
maps were obtained using a Cameca SX100 electron
probe microanalyser (EPMA) operating at 20 kV. For
observing the microstructure of the substrate alloys by
transmission electron microscopy (TEM), 3 mm discs
from the alloys were mechanically polished to a
thickness of about 100 mm. Subsequently, they were
electrochemically polished by a twin jet polisher at
235uC using 20 V. The solution for electropolishing
consisted of 68 vol.-% methanol, 10 vol.-% perchloric
4.5
3.9
3.7
4.5
4.5
4.5
4.5
W
(1.5)
(1.3)
(1.3)
(1.5)
(1.6)
(1.5)
(1.5)
4.4
4.4
4.3
2.9
3.0
3.0
3.0
(1.5)
(1.5)
(1.5)
(1.0)
(1.1)
(1.0)
(1.0)
Co
Cr
2.4 (2.5)
7.1 (7.5)
6.8 (7.5)
7.1 (7.5)
2.4 (2.6)
2.4 (2.5)
10.0 (10.5)
6.7
6.7
0.0
0.0
3.4
6.7
6.7
(8.0)
(8.0)
(0.0)
(0.0)
(4.2)
(8.0)
(8.0)
acid, 13 vol.-% ethylene glycol monobutyl ether (butyl
cellosolve) and 9 vol.-% distilled water. Since the
refractory element rich particles in the SRZ could not
be thinned by the electropolishing method, TEM foils
for observing SRZ microstructure were prepared by an
ion beam thinning technique. A Philips CM12 TEM was
used for the microstructural observation and obtaining
the electron diffraction patterns for the crystal structure
identification of the various phases. An energy dispersive
spectrometer (EDS) attached to the TEM was used to
determined the composition of the constituent phases.
Results
Substrate alloys
All the superalloy substrates developed the typical c–c9
structure in the heat treated condition. The shape of the
c9 precipitates was cuboidal in all the alloys except F-13
and F-18, where it tended to be somewhat spherical.
Since selective partitioning of refractory elements such
as Re and W to the c phase of the c–c9 substrate often
causes the precipitation of refractory rich TCP phases,5,6
the extent of enrichment of the c phase with the above
elements was determined by TEM EDS, as presented in
Table 2. The c phase compositions for the F-26, F-16
and F-13 alloys suggest that increasing the Ru content
results in a decrease in the partitioning of Re to the c
phase, which is consistent with similar observations
reported earlier.20,21 At the same time, the Ru concentration in the c phase increases. For example, while the
Re concentration of c in F-26 alloy was 6?5 at.-%, it was
3?5 at.-% in F-16 and 1?9 at.-% in F-13. The corresponding values for the Ru concentration were 6?4, 7?6
and 10?2 at.-% respectively. The same trend was also
observed between the two Cr free alloys F-13 and F-18.
The Cr content did not appear to significantly affect the
partitioning behaviour of Re or W, as evident from the
similar concentrations of these elements (about 5?0 at.%Re and 0?5 at.-%W) in the c phase of alloys F-22 and
F-20. However, in Cr free alloys (F-13 and F-18), a
much lower Re enrichment in the c phase was observed.
It is interesting to note that, in the alloys containing
Table 2 Composition of c phase of c-c9 structure of
alloys F-26, F-16 and F-13, as measured by TEMEDS technique, (at.%)
Alloy Ni
F-26
F-16
F-13
F-18
F-22
F-20
F-30
53.88
52.21
65.00
71.98
63.00
55.12
52.97
Al
Ru
Ta
Re
2.54 6.40 2.28 6.52
5.20 7.65 2.97 3.67
8.03 10.28 4.00 1.88
4.36 4.14 3.55 3.20
8.25 5.60 5.08 5.45
6.50 5.76 3.20 5.10
7.31 3.54 3.51 3.54
Materials Science and Technology
W
Co
Cr
0.98
1.10
0.23
0.68
0.52
0.68
0.37
5.64
11.55
10.58
12.09
4.40
5.08
14.73
21.76
15.65
–
–
7.70
18.56
14.03
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Formation of secondary reaction zone in single crystal superalloys
a as coated; b after 25 h cyclic exposure at 1100uC
1 Cross-sectional microstructure of b-NiAl coating on alloy F-26
relatively higher amounts of Co (.7 wt-%) such as F-16,
F-13, F-18 and F-30, the c phase was less enriched with
Re as compared to the remaining alloys.
b-NiAl coating
Cross-sectional views of the b-NiAl coating on all the
alloys revealed the typical two layer structure,22,23 as
shown in Fig. 1a for F-26. While the outer layer
consisted of single phase B2-NiAl, the inner layer,
referred to as the interdiffusion zone (IDZ), consisted of
a B2-NiAl matrix dispersed with numerous refractory
rich precipitates. The microstructural details and the
mechanism of formation of low activity b-NiAl diffusion
aluminide coatings have been widely reported.22–24 Ru
additions to Ni base single crystal alloys have been
reported to significantly restrict the formation of TCP
phases during elevated temperature exposure.25,26
However, despite high Ru contents, all the alloys
developed an SRZ layer below the coating with a typical
structure shown in Fig. 1a for F-26. The thicknesses of
the SRZs in the as coated condition for all the alloys are
provided in Table 3. The SRZ thickness ranged from
40–55 mm in all the alloys except F-16 and F-30, where it
was about half as thick (see Table 3). No SRZ was
observed in the interior of the coated samples, i.e. away
from the coatings.
As evident from Fig. 1a, the refractory rich precipitates formed in the SRZ with a needle/rod morphology.
Several precipitates are partially embedded in the IDZ,
which indicates that their precipitation initiated in the
IDZ during the aluminisation process. Subsequently,
they grew into the substrate in the form of rods as the
cellular transformation leading to the formation of the
SRZ continued. The precipitates formed in the SRZ
were rich in Re and also contained significant amounts
of W and Cr. Only in the case of F-13, the precipitates
were mostly rich in Ru and Ta. Using TEM diffraction
techniques, the Re rich phase was identified as the P
phase (Fig. 2a) and the Ru rich phase formed in F-13
alloy as b-RuAl (Fig. 2b). Unlike the P phase, b-RuAl
does not have a TCP crystal structure, but instead has a
B2, Pm3m structure. Apart from b-RuAl, a small
amount of P phase also formed in the SRZ of F-13
alloy. A typical Ru concentration profile across the
302
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coating and the SRZ, as measured using EPMA for
alloy F-26, is presented in Fig. 3. The IDZ of the coating
showed Ru enrichment in all the alloys, which can be
ascribed to the loss of Ni from this layer by outward
diffusion during the coating formation.22–24 Some
degree of outward Ru diffusion from the substrate also
occurred, as evident from the high concentration of this
element in the lower portion of the outer b-NiAl layer of
the coating (see Figs. 1a and 3).
The as coated SRZ thicknesses for alloys F-26 and F16 alloys, as mentioned in Table 3, suggest that the
higher Ru content of F-16 may be responsible in
restricting the SRZ formation in this alloy as compared
to F-26. This is consistent with the results reported by
Walston et al.,21 who observed a decreased tendency for
SRZ with increased Ru content in single crystal superalloys containing up to 3 wt-%Ru. However, it may be
noted that F-16 contained a significantly higher amount
of Co at 7?1 wt-% than F-26 (2?4 wt-%). Therefore, the
reduced SRZ growth in F-16 may also be partly due to
its higher Co content. This is supported by the fact that
a low propensity for SRZ formation was also observed
in the high Co containing alloy F-30 (10 wt-%Co).
Further, the formation of a much thicker SRZ in F-26,
which had almost the same composition as F-30 except
for the Co content, also indirectly suggests that the
higher Co concentrations was the primary reason for the
restricted cellular transformation in F-30. In the case of
F-16, however, the higher Co content is expected to be at
least partially responsible in limiting the SRZ growth.
This apparent beneficial effect of Co is consistent with
Table 3 Secondary reaction
coated alloys, mm
zone
thickness
in
b-NiAl
Alloy
As coated
After 100 h cyclic
exposure at 1100uC
F-26
F-16
F-13
F-30
F-18
F-22
F-20
41¡5
19¡3
54¡11
18¡1
45¡6
42¡2
37¡2
77¡10
22¡6
75¡6
20¡2
90¡11
64¡2
50¡4
Das et al.
Formation of secondary reaction zone in single crystal superalloys
–
a P phase formed in SRZ of F-26 alloy along with [110 ] pattern; b b-RuAl phase in SRZ of F-13 alloy along with [001]
pattern
2 Image (TEM) of TCP phases along with respective diffraction patterns
the lower degree of Re enrichment in the c phase
observed in the high Co alloys (Table 2).
The beneficial effects of Ru and Co in restricting SRZ
were not realised in alloy F-13 despite it having as high
as 14?1 wt-%Ru and 6?8 wt-%Co. Similarly, the higher
Co content of F-18 did not prove effective in restricting
the extent of SRZ (Table 2). From the SRZ thicknesses
for the alloys F-18, F-22 and F-20 (Table 3), which
contained varying Cr contents (Table 1), it appears that
the transformation kinetics are moderately inhibited by
increased Cr concentration. This is consistent with the
beneficial effect observed by Walston et al.1 in their
study on a series of Re containing superalloys where
SRZ formation decreased with increased Cr content.
Comparing F-13 and F-18, both of which did not
contain any Cr, the precipitation of b-RuAl rather than
the P phase in the transformed zone of F-13 was likely
due to its higher Ru content.
To understand the initiation and subsequent growth
of the SRZ during the coating formation in the present
alloys, a few samples of alloy F-13 were aluminised for
various durations ranging from a short period of 15 min
to longer times of 5 h. Figure 4 shows a magnified view
of the coating corresponding to 15 min aluminisation.
The presence of SRZ indicates that the cellular
3 Variation of Ru concentration across b-NiAl coating
and SRZ in alloy F-26
transformation began early during the coating process.
Further, it is evident that the b-RuAl needles of the SRZ
have actually precipitated inside the IDZ and then
grown as a part of the cellular reaction. As the
aluminisation progressed, both IDZ and SRZ grew,
although the SRZ growth was relatively much faster.
During the cyclic exposure of the coated samples at
1100uC, substantial further growth of the SRZ was
observed in all the alloys except F-16 and F-30, as
shown in Table 3. For example, in alloy F-26, SRZ
thickness increased from 41 mm in the as coated
condition to about 77 mm after 100 h exposure. The
SRZ in F-16 and F-30, however, grew very slightly from
20 to y22 mm. The SRZ microstructure, however,
remained largely unchanged, as evident in Fig. 1b for
alloy F-26 after 25 h exposure. No SRZ developed in the
interior of any of the coated samples, even after
prolonged cyclic exposure of 160 h at 1100uC.
Pt–Hf modified c–c9 coating
A typical cross-sectional microstructure of the c–c9coatings developed on all the alloys is presented in Fig. 5a
for the alloy F-26. The coating had a two layer structure
with the outer layer primarily consisting of the c9 phase.
The inner layer had a more discernable c–c9 phase
constitution and contained refractory element rich
precipitates. The c–c9 structure of the inner layer can
be more clearly seen in the coating microstructure on F30, as presented in Fig. 5b. The Pt enriched bright phase
of the inner layer in Fig. 5b is the c9 phase and that with
4 Secondary reaction zone layer developed beneath bNiAl coating in F-13 alloy after 15 min aluminisation
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5 As coated cross-sectional microstructure of Pt–Hf modified c–c9 coating on a alloy F-26 and b alloy F-30: 1
indicates outer layer and 2 indicates IDZ of coating
the grey contrast is the c phase. The phase constitution
of the two coating layers was confirmed by TEM
diffraction. Kirkendall porosity generated during the
diffusion treatment and the aluminisation process was
present in the final coating structure in several alloys
(Fig. 5a). Compared to the b-NiAl type coating
(Fig. 1a), fewer precipitates were present in the IDZ of
the c–c9 coating (Fig. 5).
Figure 6 presents the Pt and Al concentration profiles
across the c–c9 coating on F-26 alloy. For comparison
purposes, the corresponding profiles for a commercially
available b-(Ni,Pt)Al coating (MDC 150L) on superalloy CMSX-4, have been included. The Al concentration in the c–c9 coating was much lower than that in a b(Ni,Pt)Al coating. The c–c9 coating, however, had a
much higher Pt concentration than the b-(Ni,Pt)Al
coating. The measured compositions across the above c–
c9 coating can be plotted on a Ni–Al–Pt ternary phase
diagram to highlight that they indeed lie in the c9 and
6 Concentrations of Pt and Al across MDC 150L B2(Ni,Pt)Al coating on CMSX-4 substrate, and c–c9 Pt–Al
coating on alloy F-26
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7 Ni–Pt–Al ternary phase diagram for 1100uC:27 a few
composition points of coating on alloy F-26 have been
shown; dotted encircled region shows composition
range across of coating; diamonds are for as coated
and triangles are for oxidised (133 cycles at 1100uC)
coating
czc9 phase fields, as shown in Fig. 7.27 The compositions across a b-(Ni,Pt)Al coating, on the other hand,
would lie in the b phase field. Apart from having a
different structure, the c–c9 coatings also had a much
smaller thickness of approximately 25 mm than the b(Ni,Pt)Al coatings, which are typically 60–75 mm
thick.10,28 The lower thickness of the c–c9 coating is
consistent with the limited amount of Al involved in the
coating formation. No SRZ formed below the c–c9
coating in any of the alloys, as typically seen in Fig. 5
for F-26 and F-30. This is in contrast to the b-NiAl
coatings where an SRZ formed in all the alloys.
In the case of alumina forming alloys/coatings,
continued weight gain observed during oxidation usually
represents the increase in oxide scale thickness, with the
scale remaining intact on the sample surface. Once the
oxide scale starts spalling and a continuous decrease in
the sample weight is registered, the sample is no more
adequately protected against oxidation. Figure 8 presents the number of cycles at 1100uC over which positive
weight change (weight gain) was observed for the alloys
in bare and coated conditions. All the alloys in the
uncoated condition had a poor oxidation resistance,
which can be largely attributed to their relatively high
Ru contents.29 As would be expected, applying a b-NiAl
coating resulted in some degree of improvement in the
oxidation resistance. However, the c–c9 coatings provided superior oxidation resistance to the alloys. For
example, the weight gain durations for F-26, F-16 and
F-30 alloys with b-NiAl coating were 65, 64 and 60 h
respectively. The corresponding durations with the c–c9
coating were much higher at 160, 150 and 160 h
respectively. Alloys F-13, F-18, F-22 and F-20 exhibited
somewhat poorer oxidation resistance in both b-NiAl
coated and c–c9 coated conditions. The b-(Ni,Pt)Al
type Pt aluminide coatings on Ni based superalloys have
been widely reported to exhibit excellent oxidation
Das et al.
Formation of secondary reaction zone in single crystal superalloys
9 Cross-sectional microstructure of c–c9 coating on F-26
alloy after 133 cycles (100 h) of cyclic oxidation at
1100uC: phases in coating have been marked
8 Durations over which alloys showed positive weight
change (weight gain) during cyclic oxidation at 1100uC
resistance.30,31 However, it is interesting to note that the
Pt modified c–c9 coatings also impart good oxidation
resistance to superalloy substrates.
During oxidation, the coated alloys developed a
continuous alumina rich scale on the surface, as typically
shown in Fig. 9 for alloy F-26 after 100 h (133 cycles) of
exposure. The scale being predominantly alumina was
confirmed by the XRD as well as by X-ray mapping in
EPMA, as presented in Fig. 10. The above scale also
contained numerous HfO2 particles, as evident from
Fig. 10d. The presence of HfO2 in the scale was also
confirmed by XRD. The oxide layer remained adherent
for nearly 200 cycles at 1100uC for F-26, F-16 and F-30
alloys. Some degree of interdiffusion occurred between
the coating and the substrate during cyclic oxidation at
1100uC. The Pt and Al concentrations decreased
substantially as a result of such interdiffusion, as evident
from Fig. 6. The initial sharp Al concentration gradient
across the coating was eliminated after 133 cycles
(100 h) and a near constant concentration of 10 at.-%
was attained (Fig. 6). The decrease in the Al concentration was also partially caused by the scale formation.
The Pt concentration in the coating after 133 cycles
remained in the approximate range of 5 to 10 at.-%.
Despite the prolonged oxidation exposure (up to 210
cycles at 1100uC), no SRZ developed beneath the
a BSE image; b Al; c O; d Hf
10 Microstructure of c–c9 coating on alloy F-26 after 100 h cyclic oxidation at 1100uC, along with X-ray images, as
obtained using EPMA
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Formation of secondary reaction zone in single crystal superalloys
coating, as evident in Fig. 9 for alloy F-26. Owing to the
dissolution of the precipitates in the IDZ, the starting
two layer structure of the coating was no longer
identifiable after the high temperature exposure. The
coating thickness also increased appreciably during the
oxidation exposure, from 25 mm in as coated condition
to about 60 mm after 100 h exposure. Although the
coating retained its c–c9 structure throughout the
oxidation exposure, the amount of c phase increased
significantly as the coating lost Al to scale formation
(Fig. 9). This aspect is also evident from the coating
compositions presented in Fig. 7.
Discussion
The formation of a typical low activity b coating, as
those used in the present study, during aluminisation
occurs almost exclusively by outward diffusion of Ni
from the substrate. The outwardly diffused Ni reacts
with the externally provided Al to form the outer b-NiAl
layer of the coating (Fig. 1a).22,23 The loss of Ni from
the substrate leads to the formation of the inner Ni
denuded zone (IDZ), which also consists of b-NiAl
phase. Insufficient solid solubility of the alloying
elements in the in the b phase of IDZ leads to the
formation of numerous precipitates in this zone,22 as
seen in Fig. 1a.
The formation of SRZ beneath b-NiAl type aluminide
coatings is ascribed to the interdiffusion of elements
between the coating and the substrate during aluminisation.6,9 In cases where SRZ develops during the high
temperature exposure, it is suggested that the inward
diffusion of Al from the coating into the substrate causes
the transformation.9 The loss of Ru from the substrate
into the coating by outward diffusion is also believed to
be partly responsible in inducing the transformation in
Ru containing superalloys.9 In the present case of low
activity aluminisation, the coating growth occurred
primarily by the outward diffusion of Ni from the
substrate rather than inward diffusion of Al from the
coating.22,23 Therefore, it is unlikely that SRZ formation
was driven by the inward diffusion of Al from the
coating. As previously mentioned, a certain degree of Ru
loss from the substrate into the coating due to outward
diffusion was measured (Fig. 3). However, since the
starting Ru contents of the alloys were fairly high
(>5?6 wt-%), such minor loss of this element alone is not
expected to induce cellular transformation beneath the
coating during aluminisation. Apparently, the starting
c–c9 structure of the alloys was thermodynamically
unstable because of the supersaturation of the c phase
with refractory elements and Ru. In presence of the high
angle IDZ/substrate interface during the aluminisation
process, such a structure transformed to the more stable
three phase c–c9-precipitate structure, driven by the
reduction in chemical free energy resulting from
the removal of refractory elements and Ru from the
supersaturated c phase. The above transformation
occurred through a typical discontinuous precipitation
reaction where the P phase or b-RuAl precipitated from
the supersaturated c at the IDZ/substrate interface.
Subsequently, the three phases grew in a cooperative
manner with the transformation boundary gradually
moving into the untransformed substrate, as typically
observed in a cellular transformation.7,8 The IDZ
growth interface of the coating served as a high
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diffusivity path for the migration of slow diffusing
elements involved in the reaction such as Re, Ru and
Ta,5,32 and also provided the site for discontinuous
precipitation of the P/b-RuAl phase in the Ni denuded
interdiffusion zone.3,5,6 Supersaturation of Re in the c
phase was sufficiently high to cause the precipitation of
P type TCP in the SRZ of all the alloys except F-13. The
precipitation of P phase was also aided by the limited
solubility of the refractory elements in the b matrix of
the IDZ.22,23 In case of alloy F-13, the c phase was
saturated with Ru at the expense of Re (see Table 2)
which resulted in the formation of Ru rich b-RuAl
precipitates in the transformed zone.
For Re bearing superalloys containing Co in the range
12?5–20?0 wt-%, Walston et al.21 have reported that
increasing the Co content decreases the extent of the
TCP precipitation during high temperature exposure.
Although the extent of TCP precipitation and SRZ
formation may not be directly correlated,1 the present
study indicates that increasing Co does have a beneficial
effect in suppressing the extent of cellular transformation
beneath the b-NiAl coating. However, further investigation is needed to more precisely establish the effect of Co
on the SRZ formation in the above alloys. A fairly strong
effect of Cr in suppressing SRZ has been reported in a
previous study in which alloys having Cr in the range 2?3–
4?2 wt-% were examined.1 It was suggested that increased
Cr helps remove Re and W from the alloy by forming m
and s type TCP precipitates in the interdiffusion zone of
the coating, thereby reducing the chemical driving force
for the formation of P type precipitates that promote the
cellular transformation.1 As previously mentioned, only a
mild beneficial effect of Cr in suppressing the SRZ growth
was observed in the present set of alloys. However, it is
believed that the complete absence of Cr renders an alloy
extremely prone to SRZ formation, which cannot be
adequately inhibited by high Co/Ru contents, as observed
in cases of F-13 and F-18.
The cyclic oxidation at 1100uC was essentially an
extension of the aluminisation treatment in terms of the
high temperature exposure for further growth of the
SRZ. The continued SRZ growth is evident from Fig. 1b
where several TCP needles can be seen extending from
IDZ, where they precipitated during aluminisation, to
the SRZ/substrate interface, which migrated significantly inward during the cyclic exposure. Thus, the high
temperature exposure which caused oxidation damage
to the coatings also led to further SRZ growth in alloys
F-26, F-13, F-18, F-22 and F-20. However, in alloys F16 and F-30, which had restricted SRZ growth during
aluminisation, no appreciable SRZ growth occurred
during the exposure at 1100uC.
The microstructure development of the Pt–Hf modified c–c9 coating during deposition occurred in a similar
manner as the low activity b-NiAl type coatings
(Fig. 1a).22,23 However, because of extremely low Al
activity in the coating process, c9-Ni3Al rather than bNiAl formed in the outer layer of the coating. Further,
the limited loss in Ni from the substrate by outward
diffusion helped retain the c–c9 structure in the interdiffusion zone. Thus, the final coating consisted of an
outer c9-rich layer and the inner IDZ containing the
refractory rich precipitates in a c–c9 matrix. The
precipitates formed primarily in c9 phase of the IDZ
because of the relatively lower solubility of the
Das et al.
refractory elements in c9 compared to c. Further, since
the solubility is higher in c9 than in the b-NiAl, a lower
volume fraction of precipitates formed in the IDZ of the
c–c9 coating as compared to the b-NiAl coating.
Walston et al.1,6 have reported the formation of SRZ
even in the case of a Pt modified b-NiAl coating.
Therefore, it can be expected that the presence of Pt was
not the reason for which the cellular transformation was
suppressed in case of the c–c9 coatings. Gleeson et al.33
have studied interdiffusion between the CMSX-4 superalloy and the Pt containing c–c9 and b ternary alloys at
1150uC. They observed the precipitation of TCP phases
in the interdiffusion zone of the b/CMSX-4 diffusion
couple. This was in contrast to the c–c9/CMSX-4 couple,
where no such precipitation occurred. The above result
indicates that the c–c9 phase constitution is inherently
much less prone TCP precipitation as compared to the b
phase. This can be attributed to the higher solid
solubility of the refractory elements in c/c9 phases as
compared to b phase. In case of the present c–c9
coatings, some degree of TCP precipitation was
observed in the IDZ despite its c–c9 structure (Fig. 5).
However, it did not lead to cellular transformation
beneath the coating. The reason for this may be linked
to the elimination of the high angle IDZ (b)/substrate
interface, that was present in the case b-NiAl coatings.
The importance of the above interface for the occurrence of the cellular transformation can be appreciated
from the fact that the SRZ formation took place only
adjacent to the b-NiAl coating and not away from it,
even after prolonged exposure at 1100uC. Thus, it can be
concluded that the absence of a high angle interface in
case of c–c9 coatings, where both the coating and the
substrate had the same phase constitution, was the
primary reason for which the SRZ formation could be
prevented.
The superior oxidation resistance of the c–c9 coatings,
despite their low Al contents, is consistent with recent
reported results on the oxidation of ternary Ni–Pt–Al
alloys.33,34 It has been suggested that Pt addition to c–c9
compositions such as Ni–22Al–30Pt (at.-%) lowers the
oxygen solubility in the alloy. As a result, selective
oxidation of Al is promoted, ensuring the formation and
maintenance of an alumina scale even at relatively low
Al concentrations.33 Apart from the beneficial role of Pt,
the presence of Hf in the c–c9 coatings is also expected to
improve the oxidation resistance by enhancing the scale
adhesion.34
During high temperature exposure, the b-NiAl type
coatings are known to lose Al to scale formation as well
as into the substrate under the concentration gradient
that exists between the coating and the substrate.10,28
Such Al depletion from the coating causes a net volume
decrease which, along with thermal cycling, can contribute to the development of surface instabilities
(rumpling).35 The thermal mismatch stresses and the
stresses caused by the phase transformations also
contribute to the rumpling observed in case of b-type
bond coats. The phase transformations in the coating
include the b-NiAlRc9 transformation11,28,36 and the
martensitic transformation in the b phase during cyclic
heating and cooling.37 In case of the c–c9 coatings, a
shallower Al concentration gradient existed between the
coating and the substrate as compared to the b coating
(Fig. 6). Therefore, no major diffusive loss of Al from
Formation of secondary reaction zone in single crystal superalloys
the coating into the substrate is expected during
oxidation exposure. There were also no appreciable
changes in the phase constitution of the coating as it
maintained the c–c9 structure throughout the oxidation
exposure. Further, as the coating was structurally very
similar to the substrate, thermal mismatch stress would
also be comparatively much lower than that in case of btype coatings. Because of these reasons, no significant
rumpling would be expected in the c–c9 coatings even
after prolonged oxidation exposure. In the present
study, no appreciable rumpling was observed up to
200 cycles of oxidation at 1100uC (Fig. 9). Thus, based
on the results of the present study, Pt–Hf–modified c–c9
coatings appear promising as bond coats on single
crystal Ni based superalloys, particularly ‘generation 3’
and beyond alloys that have high refractory contents
and are prone to SRZ formation in presence of
traditional b-(Ni,Pt)Al coatings.
Conclusions
The microstructure, oxidation resistance and propensity
for SRZ formation were studied for a b-NiAl aluminide
coating and a Pt–Hf modified c–c9 coating on several Ru
bearing Ni based single crystal superalloys. Despite their
high Ru contents (.5?5 wt-%), all the alloys examined
in the present study developed an SRZ in presence of the
b-NiAl coating. The SRZ formation occurred via a
cellular transformation reaction, similar to that has been
observed in polycrystalline superalloys, where the c–c9
structure of the substrate transforms to a mixture of c9, c
and refractory rich precipitates. High temperature
exposure of the b-NiAl coated alloys resulted in a
significant further growth of the SRZ. In case of the c–c9
coatings, however, no SRZ formed in the superalloys,
even after prolonged high temperature exposures. These
coatings consisted of c and c9 phases which was in
contrast to the NiAl phase of the b-type coatings. The c–
c9 coatings exhibited good oxidation resistance despite
their significantly lower Al concentrations as compared
to the b-NiAl coating.
Acknowledgements
The authors wish to thank Mr C. Torbet for technical
support and Mr C. Henderson for assistance in SEM
and EPMA analysis. They also acknowledge the useful
discussion with Dr L. J. Carroll.
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