# Precipitate dissolution during deformation induced twin thickening in a
CoNi-base superalloy subject to creep

Vassili Vorontsov<EMAIL_ADDRESS>Thomas McAuliffe
<EMAIL_ADDRESS>Mark C. Hardy<EMAIL_ADDRESS>David
Dye<EMAIL_ADDRESS>Ioannis Bantounas<EMAIL_ADDRESS>Department of Design, Manufacturing and Engineering Management, University of
Strathclyde, James Weir Building, 75 Montrose Street, Glasgow, G1 1XJ, UK
Department of Materials, Imperial College, South Kensington, London SW7 2AZ,
UK  Rolls-Royce plc, Derby, DE24 8BJ, UK

###### Abstract

The tensile creep performance of a polycrystalline Co/Ni-base superalloy with
a multimodal $\upgamma^{\prime}$ distribution has been examined at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$. The rupture life of
the alloy is comparable to that of RR1000 tested under similar conditions.
Microstructural examination of the alloy after testing revealed the presence
of continuous $\upgamma^{\prime}$ precipitates and M23C6 carbides along the
grain boundaries. Intragranularly, coarsening of the secondary
$\upgamma^{\prime}$ precipitates occurred at the expense of the fine tertiary
$\upgamma^{\prime}$. Long planar deformation bands, free of
$\upgamma^{\prime}$, were also observed to traverse individual grains ending
in steps at the grain boundaries. Examination of the deformation bands
confirmed that they were microtwins. Long sections of the microtwins examined
were depleted of $\upgamma^{\prime}$ stabilising elements across their entire
width, suggesting that certain alloy compositions are susceptible to
precipitate dissolution during twin thickening. A mechanism for the
dissolution of the precipitates is suggested based on the Kolbe reordering
mechanism.

###### keywords:

Cobalt-base superalloys , Creep , Microtwinning , Atomic ordering ,
Transmission electron microscopy

††journal: arXiv.org

##  1 Introduction

The development of $\upgamma$/$\upgamma^{\prime}$ two phase Co- and CoNi-based
superalloys is promising for improved performance of high temperature
structural applications. The exceptional high temperature performance of Ni-
based superalloys is attributed to the precipitation of a coherent L12 ordered
$\upgamma^{\prime}$ phase within a disordered face centred cubic $\upgamma$
matrix. The shearing resistance offered by these ordered precipitates results
in the excellent creep resistance they display.

The discovery of a $\upgamma$/$\upgamma^{\prime}$ two phase field within the
Co-Al-W ternary system ¡1; 2¿ has enabled the development of Co- and CoNi-
based superalloys for high temperature applications ¡3; 4; 5; 6¿.
Traditionally, Ni-based superalloys are used for structural applications where
a combination of good strength, environmental resistance and creep strength
are of importance ¡7; 8¿. Understanding of creep behaviour is critical to the
implementation of Co- and CoNi-base superalloys for high performance
applications. Promising single crystal CoNi-base superalloys have been
developed with comparable creep resistance to that of 1st generation Ni-base
superalloys ¡9; 10¿.

Microtwinning has been observed in Ni-base ¡11; 12; 13; 14; 15; 16¿ and CoNi-
base ¡17¿ superalloys tested between temperatures of
$\mathrm{6}\mathrm{5}\mathrm{0}\mathrm{-}\mathrm{7}\mathrm{5}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$
and stresses of
$\mathrm{3}\mathrm{5}\mathrm{0}\mathrm{-}\mathrm{8}\mathrm{5}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$.
Though twinning is typically associated with high strain rate deformation at
low temperatures ¡18; 19¿, these microtwins form at intermediate temperatures
under slow strain rates. Knowles and Chen ¡13¿ have suggested that microtwins
nucleate from superlattice extrinsic stacking faults (SESFs).

Several mechanisms have been suggested for the creation of such faults ¡20;
21; 22; 23; 24; 13; 25¿, the reader is referred to ¡26¿ for a comprehensive
review. However, diffraction contrast experiments ¡14¿ and high resolution TEM
micrographs ¡27; 17¿ have confirmed the presence of $\frac{1}{6}$$<$$112$$>$
dislocations at the leading edge of extended stacking faults on adjacent
$\\{111\\}$ planes. Such an observation would require dislocation
decorrelation ¡28¿ and suggests that the likely twinning mechanism is that
suggested by Kolbe ¡24¿. Shearing of the L12 superlattice by
$\frac{1}{6}$$<$$112$$>$ Shockley partials results in the formation of a
pseudo twin ¡29¿. In such a configuration the order of the L12 superlattice is
not preserved across the twin boundary. Kolbe ¡24¿ suggested that a “true”
twin structure could be restored via a two step atomic re-ordering process
involving the short range diffusion of atomic species in the wake of the
leading partials.

In addition to short range atomic re-ordering, long range diffusion of
$\upgamma$ stabilising elements have been shown to segregate to the leading
partials ¡30; 31; 32¿. This Cottrell atmosphere is thought to lower the energy
penalty associated with forming a complex stacking fault (CSF) when shearing
the $\upgamma^{\prime}$ via $\frac{1}{6}$$<$$112$$>$ partials ¡33¿ and is
considered as the rate limiting step during twin formation ¡30¿. Segregation
of $\upgamma$ stabilising elements is not restricted to the leading partials
but has also been observed at the fault boundaries ¡34; 35; 36; 17¿, an
observation not limited to the Ni- and CoNi-base alloy systems but having also
been reported in MnAl Heusler alloys ¡37¿ and Mg alloys ¡38¿. In some alloy
systems, segregation of elements to stacking faults can stabilise new phases
within the fault ¡39; 35; 40; 41¿. Examples include the $\upeta$ phase in a
variant of the Ni-base superalloy ME3 ¡39¿ and the DO19 phase in the Co-Al-W-
Ta system ¡40¿.

Precipitation of carbides after prolonged high temperature exposure has been
reported in polycrystalline and single crystal Ni-base superalloys ¡42; 43;
44; 45; 46¿. In Cr rich alloys aged at intermediate temperatures, M23C6
carbides precipitate with a cube-cube orientation relationship with the parent
matrix ¡47; 45¿. Carbides are commonly associated with regions of localised
deformation and, together with grain boundaries, are thought to be dislocation
nucleation sites ¡28; 48; 49¿. In the present work, evidence suggests such
carbides to be the sources of twinning.

In this paper a $\upgamma/\upgamma^{\prime}$ CoNi-base superalloy with a
bimodal $\upgamma/\upgamma^{\prime}$ size distribution has been creep tested
to rupture at $\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$
and $\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$. Examination of
the post-test microstructure revealed the presence of precipitate-free bands
which traversed entire $\upgamma$ grains. These bands were confirmed to be
deformation-induced twins which are thought to nucleate at the incoherent
interface of grain boundary M23C6 carbides. It is postulated that the ordered
precipitates dissolve into the $\upgamma$ matrix during thickening of the
twin. A mechanism for the dissolution of the precipitates is proposed and the
consequences of these structures discussed within the text. To our knowledge,
this is the first report of precipitate dissolution during twin growth in
superalloys.

##  2 Experimental

A CoNi-base superalloy designated V208C, developed by Knop  _et al._ ¡50¿, was
melted in a vacuum arc-melter backfilled with argon and poured into a
$\mathrm{2}\mathrm{3}$×$\mathrm{2}\mathrm{3}$×$\mathrm{6}\mathrm{0}\,\mathrm{m}\mathrm{m}$
mould. Table LABEL:tab:alloy-comp shows the nominal alloy composition. The
ingot was homogenised in a vacuum furnace at
$\mathrm{1}\mathrm{2}\mathrm{5}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{4}\mathrm{8}\,\mathrm{h}$ and profile rolled at
$\mathrm{1}\mathrm{2}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ down
to a $\mathrm{1}\mathrm{3}$×$\mathrm{1}\mathrm{3}\,\mathrm{m}\mathrm{m}$ cross
cross section with a maximum of 15% strain per pass. Samples were machined
into creep specimens with a
$\mathrm{1}\mathrm{4}\,\mathrm{8}\,\mathrm{m}\mathrm{m}$ gauge length and
$\mathrm{5}\,\mathrm{m}\mathrm{m}$ gauge diameter. Annular ridges were
machined on enlarged shoulders of the test specimen for attaching an
extensometer. Creep specimens were encapsulated in quartz tubes under low
pressure argon and solution treated at
$\mathrm{1}\mathrm{1}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{1}\,\mathrm{h}$, slow cooled at
$\mathrm{2}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}\mathrm{/}\mathrm{min}$ to
$\mathrm{7}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and air cooled
(AC) to room temperature (RT). The specimen was further aged at
$\mathrm{8}\mathrm{2}\mathrm{5}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{4}\,\mathrm{h}$ followed by AC to RT.

[ caption = Nominal compositions (at. %) and $\upgamma^{\prime}$ area fraction
of alloy V208C (Patent Ref. ¡51¿)., label = tab:alloy-comp, doinside = , pos =
hbtp ]@ccccccccc—c@ Co & Ni Cr Al W Ta C B Zr $\upgamma$’ fract.
36.1341510.531 0.150.20.04 52%

The grain size of the alloy was $\mathrm{3}\mathrm{3}$ p m
$\mathrm{4}\,\mathrm{\SIUnitSymbolMicro}\mathrm{m}$, including twins, as
measured from optical micrographs using the Abrams three circle method
outlined in ASTM E112.32609. Average equivalent area diameter of the secondary
and tertiary $\upgamma^{\prime}$ precipitates were measured using ImageJ
software. The $\upgamma_{s}^{\prime}$ and $\upgamma_{t}^{\prime}$ sizes were
$\mathrm{1}\mathrm{0}\mathrm{0}\mathrm{a}\mathrm{n}\mathrm{d}\mathrm{2}\mathrm{2}\,\mathrm{n}\mathrm{m}$,
respectively.

Creep tests were performed in tension on a lever-loading machine under a
constant force to rupture111The test was paused and restarted after
$\mathrm{3}\mathrm{6}\mathrm{0}\,\mathrm{h}$ (≈$\mathrm{1}\mathrm{6}\%$
strain) as the displacement limit of the test apparatus was reached. This
involved removing the load from the specimen, cooling to room temperature and
re-adjusting the aparatus prior to restarting the test. At this point, the
material was within the tertiary creep regime with
$\mathrm{3}\mathrm{0}\,\mathrm{h}$ remaining to rupture.. The test temperature
was set to $\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
the starting stress to
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$. Two thermocouples
attached with high temperature wire to the centre of the gauge were used to
monitor temperature during testing. Creep strain was recorded by fitting an
extensometer to the annular ridges on the test piece. The change in
displacement was measured using a transducer attached to a pair of arms
extending below the furnace.

The deformed microstructure was examined by sectioning the gauge perpendicular
to the loading direction. Samples were progressively ground down to 4000 grit
SiC paper and then polished using a suspension containing 100ml water, 30ml
coloidal silica suspension and 70ml H2O2. The $\upgamma^{\prime}$ precipitates
were revealed by electroetching in a solution of 2.5% phosphoric acid in
methanol using a voltage of 10V for 10 seconds. Transmission electron
microscopy (TEM) lamellae were prepared from the polished cross-section using
an FEI Helios Nanolab 600 focussed ion beam workstation via the _in-situ_
lift-out technique ¡52; 53¿. The twin chemistry was examined in a JEOL
JEM-2100F microscope equipped with a XEDS silicon-drift detector from Oxford
Instruments. XEDS data was collected and analysed using the Oxford Instruments
AZTec v3.1 software package. Electron micrographs were acquired with the
InLens detector on Zeiss Sigma 300 and a Zeiss Auriga Cross-Beam scanning
electron microscopes. The secondary electron images presented were acquired on
a FEI Helios Nanolab 600 workstation.

##  3 Results

###  3.1 Creep Behaviour

The tensile creep strain vs time curve for the alloy tested at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$ is shown in Figure 1
(a). The time to rupture was $\mathrm{3}\mathrm{9}\mathrm{0}\,\mathrm{h}$ with
a minimum creep rate of
$\mathrm{8}\,\mathrm{6}\mathrm{E}\mathrm{-}\mathrm{9}\,\reciprocal\mathrm{s}$.
Figure 1 (b) shows the test results displayed on a Larson-Miller parameter
(LMP) plot, equation 1.

$\text{LMP}=T\cdot(20+\log_{10}t_{r})\times 10^{-3}\,,$ (1)

where $T$ is the test temperature in Kelvin and $t_{r}$ the time to rupture in
hours.

The graph compares the current alloy to the performance of industrially
relevant disk alloys U720Li ¡54¿, RR1000 ¡55¿ and ME3 (P/M) ¡56¿. At a stress
of $\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$ V208C in the
current microstructural condition has a comparable LMP to that of RR1000 ¡55¿,
a disk superalloy currently used in gas turbine engines.

Figure 1: (a) Creep curve of specimen tested in tension at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$, and (b) comparison of
creep performance with U720Li ¡54¿, RR1000 ¡55¿ and ME3 ¡56¿ on a Larson
Miller parameter graph.

###  3.2 Grain Boundary Evolution

Exposing the alloy to a temperature of
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{3}\mathrm{9}\mathrm{0}\,\mathrm{h}$ precipitated out Cr rich M23C6
carbides and $\upgamma^{\prime}$ at the $\upgamma$ grain boundaries. The grain
boundary structure before creep testing is shown in Figure 2 (a). A dashed
line is drawn on the micrograph to outline the boundary, as there is little
contrast to distinguish it from the bulk of the two adjacent grains. Few
carbides can be seen decorating the boundary. The post creep test
microstructure is shown in Figures 2 (b) and (c). Micrographs from the gauge
section after exposure at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$ for
$\mathrm{3}\mathrm{9}\mathrm{0}\,\mathrm{h}$ reveal the presence of phases
decorating the grain boundaries of the alloy, Figure 2 (b). These phases are
also observed within the microstructure from the threads of the specimen,
Figure 2 (c), implying that they are not stress induced phases but rather a
result of heat treatment during testing.

Figure 2: InLens micrographs showing (a) grain boundary carbides in the pre-
test microstructure, (b) and (c) growth of continuous grain boundary phases
post creep test at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{3}\mathrm{9}\mathrm{0}\,\mathrm{h}$ in the specimen gauge and thread,
respectively.

The grain boundary phases which formed after testing are Cr rich M23C6
carbides and the L12 $\upgamma^{\prime}$ phase. Figure 3 (a) shows a centre
beam dark field (CBDF) conventional (C)TEM micrograph of a foil prepared from
the gauge of the creep specimen. Two grains labelled G1 and G2 are visible in
the top left and bottom right of the micrograph, respectively. Grain G2 is
oriented along the $[001]$ zone axis, Figure 3 (c), with a superlattice
reflection selected for the CBDF micrograph. The $\upgamma^{\prime}$
precipitates appear bright against a dark $\upgamma$ matrix, while grain G1
does not contribute to the image and appears dark. A carbide is located along
the grain boundary, more clearly visible in the centre beam dark field (CBDF)
micrograph in Figure 3 (b). The carbide is oriented along the $[001]$ zone
axis and its crystal structure resembles that of M23C6, Figure 3 (d).

Scanning transmission electron microscopy (STEM) energy dispersive
spectroscopy (EDS) maps reveal that the carbide from Figure 3 (b) is enriched
in Cr, Figure 3 (f). Additionally, the carbide is surrounded by
$\upgamma^{\prime}$ precipitates and bound between two Ta/Zr rich carbides.
The compositions of the Cr and Ta/Zr-rich carbides are presented in Table
LABEL:tab:carbide-comp.

[ caption = Carbide compositions (at. %) measured via STEM-EDS., label =
tab:carbide-comp, doinside = , pos = hbtp ]@lrrrrrrrr@ Phase & Co Ni Al W Ta
Cr C Zr M23C6 6 2 0 3 0 75 13 0 MC 5 2 0 4 24 31 26 9

The Cr enrichment coupled with the large cubic lattice parameter suggests that
this is an M23C6 carbide with approximate composition Cr23C6. The space group
of this phase is classified as $Fm\mathrm{bar}3m$ in Hermann-Mauguin notation.
In turn, Ta/Zr-rich carbides tend to be MC carbides. The formation of M23C6
and $\upgamma^{\prime}$ at the grain boundaries after prolonged exposure to
high temperature is due to the decomposition of MC grain boundary carbides
according to the
$\text{MC}+\upgamma\rightarrow\text{\ce{M23C6}}+\upgamma^{\prime}$ reaction
sequence ¡7; 8¿.

###  3.3 Carbide Induced Slip

Our results show that the growth of M23C6 follows a cube-cube relationship
with one of the grains while inducing strain and the nucleation of slip bands
in the adjacent grain, Figure 4. A selected area diffraction pattern (SADP) of
grain G2 imaged along the $[001]_{\upgamma}$ zone is shown in Figure 3 (c), a
SADP of the M23C6 carbide which shares a cube-cube relationship with this
grain is shown in Figure 3 (d). A SADP of both the carbide and $\upgamma$
grain are presented in Figure 3 (e). These diffraction patterns clearly
display the orientation relationship between the carbide and matrix. A CBDF
micrograph of the carbide which shares a relationship with grain G2 is
provided in Figure 3 (b). This carbide is labelled as “M23C6||G2” in Figure 4
(a) and (c), denoting it’s cube-cube orientation relationship with grain G2,
and is connected to a deformation twin in the adjacent grain G1. Likewise,
carbide “M23C6||G1” shares a cube-cube relationship with grain G1 and is
connected to a deformation band in the adjacent grain G2.

We can infer that the carbides nucleate and form a semi-coherent interface
with the grain they share an orientation relationship with while exerting a
strain on the adjacent grain with which they share an incoherent interface
during growth. Coupled with the load from the creep test, these incoherent
interfaces make for favourable sites for the initiation of slip events.
Delaying the precipitation of these carbides could thus prolong the creep life
of the alloy.

Figure 3: CBDF micrographs from grain G2 showing (a) intragrannular and grain
boundary $\upgamma^{\prime}$ precipitates and (b) grain boundary M23C6 carbide
coherent with grain G2. SADPs from (c) the bulk of grain G2 and (d) the M23C6
carbide. (e) SADP showing the cube-cube orientation relationship between G2
and the M23C6 carbide. All SADPs imaged along the $[001]$ beam direction. (f)
STEM HAADF micrograph of carbide with accompanying EDS maps. All micrographs
acquired from gauge of creep specimen tested at
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$ and
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ for
$\mathrm{3}\mathrm{9}\mathrm{0}\,\mathrm{h}$. Figure 4: (a) BF TEM micrograph
taken along the $<$$100$$>$ zone axis of grain G1 and carbide M23C6||G1
showing deformation twin intersect the foil plane along the $<$$110$$>$ trace.
(b) SADP from (a) - Note that the vertical streaks in the DP are due to
charging of the CCD camera and do not reflect a property of the material. (c)
BF TEM micrograph showing deformation twin in G1 ending at grain boundary
carbide labelled M23C6||G2 which shares a cube-cube relationship with grain G2
and an incoherent interface with G1. (d) Schematic showing defect nucleation
at an incoherent interface as a result of grain boundary M23C6 carbide
precipitation and growth.

###  3.4 Microstructural Evolution

A bimodal $\upgamma^{\prime}$ microstructure consisting of coarse secondary
and finer tertiary precipitates was obtained after heat treating the rolled
ingot. The microstructure is presented in Figure 5 (a). It consists of
secondary $\upgamma^{\prime}$ precipitates with a mean diameter of
$\mathrm{1}\mathrm{0}\mathrm{0}\,\mathrm{n}\mathrm{m}$ and
$\mathrm{2}\mathrm{2}\,\mathrm{n}\mathrm{m}$ tertiary precipitates. The
seconary and tertiary precipitate area fractions were 49 and 3%, respectively.
After creep testing the secondary $\upgamma^{\prime}$ precipitates coarsen and
coalesce at the expense of the tertiary, Figure 5 (b). In addition, bands free
of precipitates which traverse entire $\upgamma$ grains are observed. These
bands, later shown to be deformation microtwins, are the focus of subsequent
sections.

Figure 5: InLens micrographs of (a) the as-heat treated microstructure showing
the secondary and tertiary $\upgamma^{\prime}$ precipitates, and (b) post-test
microstructure from the gauge showing coarsening and coalescence of secondary
$\upgamma^{\prime}$ along with precipitate-free bands within the grain
interior.

###  3.5 Deformation Microstructure

Figure 6 (a) shows a section of the microstructure of the crept specimen taken
from the gauge perpendicular to the loading direction. The arrows indicate
precipitate free bands. These bands end at grain boundary steps. Such steps
are responsible for the nucleation of grain boundary cavities leading to
tertiary creep. Grain boundary $\upgamma^{\prime}$ and Cr rich M23C6 carbides,
characterised in the previous section, can also be seen to decorate the grain
boundary.

Figure 6: Secondary electron micrographs showing (a) precipitate free bands
(white arrows) terminating at a grain boundary, (b) magnified view of
precipitate free bands from the grain interior, and (c) location of TEM foil
encompassing multiple bands.

A magnified view of the grain interior showing a deformation band is displayed
in Figure 6 (b). The figure highlights the abscence of precipitates within the
band. In order to characterise the deformation and chemical structure of the
band, a TEM foil was extracted from a deformed grain. Figure 6 (c) shows the
location of the TEM lamella prior to FIB milling. The area was selected in
order to encompass multiple bands within the foil.

Conventional transmission electron microscopy reveals that these deformation
bands are deformation twins on the order of a few hundred nanometers in
thickness. A selected area diffraction pattern (SADP) from an area containing
the precipitate free bands is shown in Figure 7 (a). The foil is oriented such
that the beam direction is along the $[011]$ zone axis of the grain. Along
this zone, diffuse diffraction spots can be seen about the
$(\mathrm{bar}1\mathrm{bar}11)$ plane, revealing the twin nature of the
precipitate free bands.

Figure 7: (a) Selected area diffraction pattern along the $[110]$ zone showing
twin reflections (marked with a ‘T’ subscript), and centre beam dark field
micrographs from (b) matrix $[\mathrm{bar}200]$ beam with twins appearing dark
and (c) twin $[1\mathrm{bar}11]_{T}$ reflection with twins appearing bright.

A $[\mathrm{bar}200]$ matrix reflection was selected to create a centre beam
dark field (CBDF) micrograph in Figure 7 (b). The parent grain is bright and
the deformation bands dark. Figure 7 is a CBDF micrograph using a
$[1\mathrm{bar}11]$ twin reflection. In this instance the parent grain appears
dark and the twin structure bright. These micrographs confirm that the
deformation bands observed in the SEM are deformation twins. Consistent with
the absence of $\upgamma^{\prime}$ precipitates observed in the SEM
micrographs, no superlattice reflections can be observed from the twinned
reflections. Despite the lack of superlattice reflections the chemical
compsition of the twin suggests regions of high concentration of
$\upgamma^{\prime}$ stabilising elements.

###  3.6 Twin Chemistry

Consistent with the majority of observed deformation twins from SEM
micrographs the twin channels are enriched in $\upgamma$ stabilising elements,
as shown via STEM-XEDS maps in Figure 8(b). The X-ray maps show the
distribution of Co, Cr, Ni and Al within two twinned regions. In this alloy,
Co and Cr partition to the $\upgamma$ phase while Ni and Al partition to the
$\upgamma^{\prime}$ phase ¡50¿. Despite the SEM micrographs revealing an
absence of $\upgamma^{\prime}$ precipitates within the deformation channels,
the elemental maps show concentrations of $\upgamma^{\prime}$ stabilising
elements coupled with regions rich in $\upgamma$ stabilising elements. The
composition of elements across the twin in both such cases is examined next.

Line profiles across the thickness of the twin are presented from regions rich
in $\upgamma^{\prime}$ and $\upgamma$ stabilising elements, Figure 8(a) and
(b), respectively. We first examine the case of a twinned region containing an
$\upgamma^{\prime}$ precipitate, or more specifically, a region rich in Ni and
Al $\upgamma^{\prime}$ stabilising elements, Figure 8(a). Tracking the Ni Kα
quantitative line profile from left to right shows a decrease in concentration
to a minimum of ≈$\mathrm{2}\mathrm{4}\,\mathrm{a}\mathrm{t}\,\%$ followed by
a rise to ≈$\mathrm{3}\mathrm{6}\,\mathrm{a}\mathrm{t}\,\%$. This first dip is
believed to be due to the converged electron beam crossing from a
$\upgamma+\upgamma^{\prime}$ region of the foil to a $\upgamma$ channel, the
composition of which is consistent with atom probe tomography results of the
$\upgamma$ phase in this alloy ¡50¿. The beam then scans across a
$\upgamma+\upgamma^{\prime}$ region of the foil, which becomes progressively
richer in $\upgamma^{\prime}$. Following the peak in Ni composition of
≈$\mathrm{3}\mathrm{6}\,\mathrm{a}\mathrm{t}\,\%$ there is a dip in
concentration of ≈$\mathrm{6}\,\mathrm{a}\mathrm{t}\,\%$ Ni at the twin/matrix
interface. The Ni concentration increases throughout the thickness of the twin
to ≈$\mathrm{3}\mathrm{4}\,\mathrm{a}\mathrm{t}\,\%$, however does not recover
to the composition of the undeformed precipitate. A second dip in the
concentration is again observed at the adjacent twin/matrix interface. After
the electron beam crosses this planar interface the Ni concentration gradually
decreases from a $\upgamma+\upgamma^{\prime}$ phase region to that of pure
$\upgamma$. The Al Kα line profile follows a similar trend to that of Ni, both
elements preferentially partitioning to the $\upgamma^{\prime}$ phase.

Figure 8: X-ray line profiles across thickness of twin showing (a) enrichment
of $\upgamma$ and depletion of $\upgamma^{\prime}$ stabilising elments at
twin/matrix interface when the twin is confined within a $\upgamma^{\prime}$
precipitate, (b) consumption of $\upgamma^{\prime}$ precipitates by $\upgamma$
stabilising elements during twin thickening. Schematics (c) and (d) show the
physical representations of the line profile observations from (a) and (c),
respectively. Hashed lines represent the ($\mathrm{bar}101$) planes. ‘P’ and
‘T’ in the schematics denote regions of the parent and twin, respectively.

Line profiles of X-ray peaks from the $\upgamma$ stabilisers, Co Kα and Cr Kα,
follow an inverse pattern to that of Ni and Al. The concentration of Co and Cr
gradually decreases as the electron beam moves from the $\upgamma$ channel to
a region containing $\upgamma+\upgamma^{\prime}$ phases. At the twin/matrix
interfaces the concentration of $\upgamma$ stabilising elements increases. A
schematic of the precipitate and twin configuration is depicted in Figure 8(c)
where a twin is confined within a $\upgamma^{\prime}$ precipitate with its
twin/matrix interfaces depleted in $\upgamma^{\prime}$ and enriched in
$\upgamma$ stabilising elements. This is attributed to long range diffusion of
$\upgamma$ stabilising elements during twin lengthening through a
$\upgamma^{\prime}$ precipitate ¡31; 32¿ and helps to lower the energy penalty
associated with the passage of $\frac{1}{6}$$<$$112$$>$ partials on adjacent
$\\{111\\}$ planes during twin thickening ¡17¿.

A different scenario is presented in Figure 8(b) where the composition across
the entire thickness of the twin resembles that of the $\upgamma$ phase.
Tracking the Ni Kα line profile. The concentration of Ni gradually increases
from ≈$\mathrm{2}\mathrm{6}\,\mathrm{a}\mathrm{t}\,\%$, in the $\upgamma$
phase, to ≈$\mathrm{3}\mathrm{9}\,\mathrm{a}\mathrm{t}\,\%$ as it enters a
$\upgamma+\upgamma^{\prime}$ phase region. A sudden drop to
≈$\mathrm{2}\mathrm{6}\,\mathrm{a}\mathrm{t}\,\%$ is observed at the
twin/matrix interface. Up to this point, this observation is consistent with
the previous case when the twin is contained within a $\upgamma^{\prime}$
precipitate. However, unlike the previous example, the concentration of Ni
does not recover to that of the $\upgamma^{\prime}$ precipitate but remains
low at ≈$\mathrm{2}\mathrm{5}\,\mathrm{a}\mathrm{t}\,\%$, resembling that of
the $\upgamma$ phase, throughout the thickness of the twin. As the converged
beam crosses the adjacent twin/matrix interface the concentration of Ni
sharply rises again by ≈$\mathrm{5}\,\mathrm{a}\mathrm{t}\,\%$ and then
follows a gradual increase to
≈$\mathrm{3}\mathrm{4}\,\mathrm{a}\mathrm{t}\,\%$ before gradually decreasing
to ≈$\mathrm{2}\mathrm{7}\,\mathrm{a}\mathrm{t}\,\%$. The raised Ni
concentration on either side of the twin is thought to be two separate
$\upgamma^{\prime}$ precipitates. The Al Kα peak follows a similar trend to
that of Ni, while the Co and Cr Kα peaks, being $\upgamma$ stabilisers, follow
an inverse pattern. I.e., deplete when the electron beam scans the
$\upgamma^{\prime}$ precipitate and increase in the twin region.

In this instance we believe that the twin has nucleated (i.e. lengthened)
within the $\upgamma$ matrix and consumed the adjacent $\upgamma^{\prime}$
precipitates during the thickening stage of its life. A schematic is shown in
Figure 8(d) to depict the migration of $\upgamma$ stabilising elements to the
prior precipitate boundary. This creates a region along the length of the twin
depleted in $\upgamma^{\prime}$ stabilising elements with a composition very
close to that of the $\upgamma$ phase which stretches across the entirity of
its width. A mechanism is proposed in the disussion for the creation of such a
twinned region.

##  4 Discussion

###  4.1 Carbide precipitation and slip

The precipitation of M23C6 carbides at the grain boundaries introduces strain
in the adjacent grain which, coupled with the load from the creep test, make
the carbides favourable areas for slip nucleation. The work from Section 3.2
clearly shows grain boundary M23C6 carbides precipitating coherently with one
of the grains adjacent to the grain boundary. Figure 4(c) shows two M23C6
carbides precipitating coherently with one of the grains and inducing slip on
the adjacent grain. Carbide M23C6||G1 shares a cube-cube orientation
relationship with grain G1. The carbide shares a semi-coherent interface with
this grain while its interface with the adjacent grain G2 is incoherent. A
deformation twin is seen to lie on the adjacent grain G2 intersecting grain G1
at the incoherent interface of the M23C6||G1 carbide. The same can be observed
for carbide M23C6||G2 which has a cube-cube orientation relationship with G2
and induces slip in the adjacent grain G1. A schematic depicting these
observations is presented in Figure 4(d).

In support of the above argument, SEM micrographs from the threads of the
creep specimen reveal that stress is not a prerequisite for the formation of
these grain boundary carbides, Figure 2 (c). This means that in trying to
understand the spacial coincidence of the carbide and deformation structure we
are lead to believe that the defects are not a prerequisite for carbide
nucleation but that likely the opposite is true. Carbides precipitate at the
grain boundaries due to thermodynamics and act as nucleation sites for the
deformation structures.

The implication of this observation is that there could be room to improve the
creep performance of this alloy by lowering the Cr additions. Grain boundary
M23C6 is stabilised by chromium. Reducing chromium could suppress the
formation of the grain boundary carbide during exposure to high temperature.
If the carbides indeed act as slip initiation sites for twin formation, as
discussed above, their suppression could also delay the onset of such
deformation. This in turn would prolong the time to rupture and improve the
creep performance of the alloy. Given the comparable performance of this alloy
with RR1000 ¡55¿, a Ni-base superalloy in current use in jet engines, there is
therefore potential for polycrystalline CoNi-base superalloys to outperform
current generation disk Ni-base superalloys.

###  4.2 Twinning mechanism

In this study we are concerned with twin growth and how it relates to creep
rupture. We observe long regions of precipitate-free twins which give rise to
grain boundary steps, believed to be the cause for the transition to the
tertiary creep regime and subsequently rupture. Two mechanisms are proposed
for twin growth based on the presence of $\upgamma^{\prime}$ precipitates
within the vicinity of the original twin. Both cases can be accounted for by
the following example.

A single twin is nucleated from a SF that has grown across both the $\upgamma$
and $\upgamma^{\prime}$ phases from one end of a grain to the other. Across
the length of the fault both phases have undergone a displacement. As the twin
grows perpendicular to the $\\{$$111$$\\}$ twin plane it carries with it the
segregation of $\upgamma$ stabilising elements. When the twin length is
confined to one $\upgamma^{\prime}$ precipitate, Figure 8(c), the twin
boundaries are enriched in Co and Cr and depleted in Ni and Al. In instances
where the twin has nucleated within a precipitate and continues to grow into
the $\upgamma$ matrix, once the twin boundary reaches the adjacent precipitate
instead of cutting into the precipitate, dissolution of the precipitate
occurs, Figure 8(d). The twin boundary consumes the neighbouring precipitate.

###  4.3 Proposed precipitate dissolution mechanism

As the twin boundary moves by the motion of $\frac{1}{3}$$<$$112$$>$ partials
gliding on adjacent $\\{$$111$$\\}$ planes, the twin thickens one atomic plane
at a time, Figure 9(a). If the twin is locally contained within a region of
disordered $\upgamma$ phase there will come a time during the thickening
process were the twin boundary will interact with a
$\upgamma/\upgamma^{\prime}$ interface, Figure 9(b). The
$\frac{1}{3}$$<$$112$$>$ partial is a $\frac{1}{6}$$<$$112$$>$ superlattice
partial in $\upgamma^{\prime}$. Shearing of the first $\\{111\\}$
$\upgamma^{\prime}$ atomic layer by the superlattice partial creates a single
layer pseudo-twin within the $\upgamma^{\prime}$ precipitate. If this pseudo-
twin was in the bulk of the precipitate atomic reordering with atoms from an
adjacent $\\{111\\}$ plane would convert it to a true twin structure. However,
in the current case we are presented with a single layer pseudo-twin boundary
on one side of the planar fault and a $\upgamma/\upgamma^{\prime}$ interface
at the other, Figure 9(c). We hypothesise that such a configuration is
thermodynamically unstable at the testing temperature of
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and so the
single layer pseudo-twin would disorder, via atomic reshuffling with the
$\upgamma$ phase, dissolving into the $\upgamma$ matrix, Figure 9(d). This
results in a $\upgamma/\upgamma^{\prime}$ twin boundary as observed in the
present work.

Figure 9: Schematic showing formation of a single layer pseudo twin at a
$\upgamma-\upgamma^{\prime}$ interface and consequent precipitate dissolution.
(a) Twin within the $\upgamma$ matrix, dark grey region. (b) Twin growth into
the first $\upgamma^{\prime}$ atomic layer showing creation of a single layer
pseudo twin in the wake of the leading Schockley partial, light grey region.
(c) The partial has made its way across a narrow stretch of the
$\upgamma-\upgamma^{\prime}$ interface leaving behind it a single layer pseudo
twin. A parent-twin $\upgamma^{\prime}$-psuedo twin interface is shown on the
left and a pseudo twin matrix on the right. (d) Atomic shuffling in the wake
of the leading Shockley partial disorders the single layer pseudo twin slowly
dissolving the $\upgamma^{\prime}$ precipitate during the process of twin
growth.

Such a mechanism is important for two reasons. (i) The layer-by-layer
dissolution process of the ordered precipitate phase at the
$\upgamma/\upgamma^{\prime}$ twin interface may be a rate limiting step for
twin thickening. (ii) The formation of precipitate free bands which traverse
entire $\upgamma$ grains can provide resistance-free dislocation paths, which,
in turn increase creep rates and result in large strains at grain boundaries.
This may be more prevalent during a transition to a high stress low
temperature creep regime where twin motion is not favoured and dislocations
glide free from the resistance of ordered precipitates along previously formed
twins.

##  5 Summary and conclusions

The tensile creep performance of a polycrystalline Co/Ni-base superalloy
processed so as to produce a multimodal $\upgamma^{\prime}$ distribution has
been examined at $\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$
and $\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$. The main findings
are as follows:

  * 1.

The rupture life of the alloy is similar to that of commercial alloy RR1000
tested under the same conditions.

  * 2.

M23C6 carbides nucleate at $\upgamma$ grain boundaries during creep. These
carbides observe a cube-cube orientation relationship with the grain which
they share a semi-coherent interface with, while acting as defect nucleation
sites in the adjacent grain with which they share an incoherent interface.

  * 3.

$\upgamma^{\prime}$ precipitate-free twins are observed to cross large regions
of $\upgamma$ grains. Solute segregation at the $\upgamma$/$\upgamma^{\prime}$
phase boundary is suggested to result in the dissolution of the precipitate as
the twin grows from the dissordered matrix phase into the ordered precipitate,
slowly consuming the precipitate and creating a twin predominantly free of the
ordered $\upgamma^{\prime}$ phase. This mechanisms is similar to the Kolbe re-
ordering with the exception that in the classical mechanism the pseudo twin is
contained within the ordered precipitate whereas in the present case it is
sandwiched between the ordered precipitate and the disordered matrix phase.

These findings suggest that polycrystalline Co/Ni-base superalloys have the
potential to compete with commercial Ni-base superalloys in terms of creep
performance. Precipitation of M23C6 carbides act as defect nucleation sites at
the incoherent carbide-matrix interface, and that the composition of the V208C
Co/Ni-base superalloy is prone to precipitate dissolution during twin growth
when creep tested in tension at
$\mathrm{8}\mathrm{0}\mathrm{0}\mathrm{\SIUnitSymbolCelsius}$ and
$\mathrm{3}\mathrm{0}\mathrm{0}\,\mathrm{M}\mathrm{Pa}$.

## Acknowledgement

This work was funded under the Rolls-Royce-EPSRC strategic partnership in
structural metallic systems for gas turbines (EP/M005607/1). The authors would
like to thank Dr Catrin Davies and Alex Toth from the Mechanical Engineering
Department at Imperial College London for help with the use of the creep rigs
within the High Temperature Testing laboratory. Our thanks is also extended to
Mike Lennon and Ben Wood for their assistance in creep specimen preparation.
VAV would like to acknowledge support from Rolls-Royce plc and Imperial
College London under the Imperial College Research Fellowship scheme.

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