Creep behavior of a γ′-strengthened Co-base alloy with zero γ/γ′-lattice misfit at 800 °C, 196 MPa [original]

ARTICLE
Creep behavior of a c 9 -strengthened Co-base alloy with zero
c / c 9 -lattice mis ﬁ t at 800 °C, 196 MPa
Jan Midtlyn g and Alexande r I. Epishin
a)
Institute of Materials Science and Technologies, Technical University of Berlin, Berlin 10587, Germany
Nikolay V. Petrus hin
All-Russian Scienti ﬁ c Research Institute of Aviation Materials (VIAM), Department of Superalloys, Moscow
105005, Russia
Thomas Link
Institute of Materials Science and Technologies, Technical University of Berlin, Berlin 10587, Germany
Gert Nolze
Federal Institute for Materials Research and Testing (BAM), Department of Materials Engineering, Berlin 12205,
Germany
Igor L. Svetlov
All-Russian Scienti ﬁ c Research Institute of Aviation Materials (VIAM), Department of Superalloys, Moscow
105005, Russia
Walter Reime rs
Institute of Materials Science and Technologies, Technical University of Berlin, Berlin 10587, Germany
(Received 28 April 2017; accepted 11 October 2017)
Deformation and structural behavior of an expe rimental c 9 -strengthened Co-base alloy during creep at
800 °C and 196 MPa have been investigated. The ch aracteristic features of this alloy are zero
c / c 9 -lattice mis ﬁ ta n da ﬁ ne c / c 9 -microstructure. In th e initi al conditi on, the c 9 -precipitates in this
alloy are small (size of about 100 nm), have polyhedral morphology, and are separated by the very
narrow c -channels (width of a bout 10 n m). The tests performed up to about 1% creep strain (about
500 h creep time) gave creep curves with a slow constant strain rate and without an a pparent transient
creep, typi cal for sup eralloys wi th nonzero mis ﬁ t. In this initial stage of creep, ente ring of the narrow
c -channels by dislocations is blocked by a strong Orowan force. The micromechanism of creep was
identi ﬁ ed as an octahedral glide of h 011 i superdislocations simultaneously in two phases, c and c 9 .
The c / c 9 -microstructure with zero mis ﬁ t shows no rafting bu t rapidly coarsens isotropically. It is
concluded that zero mis ﬁ ti sb e n e ﬁ cial at the initial stages of the c reep but is unfavourable fo r long-
term creep because of the continuous microstructural coarsening.
I. INTRODUCT ION
Recently, Co-base alloys with a c / c 9 -microstructure
similar to that of Ni-base superalloys were discovered.
1
In these Co-base alloys, a face-centred cubic (fcc) solid
solution of Co, c , is strengthened by precipitation of
a stable ternary Co
3
(Al,W) intermetallic compound, c 9 ,
with the ordered structure L1
2
(Cu
3
Au type). It is
assumed that these new c 9 -strengthened Co-base alloys
can substitute the Ni-base superalloys in some important
applications. The reasons for this assumption are:
(i) The discovered Co-base alloys have higher solidus
and liquidus temperatures compared with those of Ni-
base superalloys, which makes them potentially attractive
for high-temperature applications.
(ii) These Co-base alloys have good castability due to
the narrow solidi ﬁ cation interval and a low degree of
segregation during solidi ﬁ cation. This is important for
manufacturing of the large single-crystal blades needed
for power gas turbines.
(iii) These Co-base alloys are c single phase in a wide
temperature interval, and the material is then soft and
ductile. This could permit high-temperature processing,
e.g., forging or rolling, assuming that these alloys could
be developed as wrought or sheet materials.
Research activities on these new Co-based alloys
include: alloy development,
1 – 7
mechanical testing,
2 – 8
testing of structural stability
9,10
and oxidation,
2,11
and
investigations of deformation micromechanisms.
3,12 – 15
At the current stage of alloy development, several
Co-base alloys have been proposed, which have creep
strengths at temperatures 800 – 900 °C approaching those
of Ni-base superalloys, e.g., Refs. 3 – 5.
Contributing Editor: Gunther Eggeler
a)
Address all correspondence to this author.
e-mail: [email protected]
DOI: 10.1557/jmr.2017.424
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 Ó Materials Research Society 2017 4466
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An in te re s ti ng fe at ure of ne w c 9 -str engt he ned Co -ba se
al lo ys is that th e c / c 9 -lat ti ce mi s ﬁ t in th es e al lo ys
is sig n i ﬁ ca nt ly sh if te d to wa rd th e po s it iv e si d e. Fo r
exa mp le , in the exp erim enta l allo y Co9 Al9W0 .1 B, the
“ co nstr ai ned ” c / c 9 -m is ﬁ t (i. e., m is ﬁ t bet we en th e sp acin gs
of co nstr ain ed c -a n d c 9 -l atti ces mea sur ed in unde for med
ma teri al) was fo und to be 1 0. 8%.
12
Fo r comp aris on , the
mis ﬁ t in the wid ely used s upe ral loy CM SX-4 is of ab out
 0. 25%.
16
Su ch a bi g diff ere nce be twee n th e mis ﬁ ti nt h e
Co- an d in Ni -b as e allo ys res ult s from a d iffe ren ce in th e
par titi oni ng of the el em ents bet wee n the c -a n d c 9 -p has es.
As fol low s from , e. g.,
1 – 7
el emen ts like Mo , W, Re,
al thoug h bein g th e matri x ele men ts in Ni -b ase al loys , in
Co -ba se all oy s have a pr ef er en ce t o part it ion i nt o t he
c 9 -p hase , by th is in crea sin g its la tt ice sp acin g mor e th an in
t he N i- ba se . T he ﬁ n din g of a po sitiv e mis ﬁ t in Co- bas e
al loys r eop ens th e dis cus sion ab out th e imp or tanc e of sig n
and ma gni tude of c / c 9 latt ice mi s ﬁ t in sup eral lo ys, see
Mu ghrab i.
17
Clari ﬁ ca tion o f th is qu esti on will ne ed
th eo reti cal in ves tig atio ns, e. g. , sim ilar to R ef. 1 8, as we ll
as th e mis ﬁ t- rele van t expe ri men ts. Th is arti cle p resen ts
su ch an exp erim enta l con trib uti on.
The alloy under investigation in this work is
a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice
mis ﬁ t that was developed recently at the VIAM, Mos-
cow. The points of interest are the deformation behavior,
structural behavior, and deformation mechanism of this
superalloy at the initial stage of creep at 800 °C under
a load of 196 MPa.
II. EXPERIMENTA L
The alloy under investigatio n was an experimental
Co-base alloy VIAM-M2 (M: Multicomponent) developed
at the VIAM in Moscow, with composition as given at the
top of Table I. It is seen from Table I that this alloy contains
a high concentration of Co, 53.1 at.%, and the sum of the
concentrations of Co and a nother basic element, Ni ( 15.9 at.
%), is 69 at.%, which is t ypical for Ni-ba se superalloys.
Concentrations of other alloying elements are close to those
in Ni-base superalloys. In sp ite of the compositional and
structural ( c / c 9 ) similarit y with Ni-base alloys, the element
partitioning in the Co-base alloy essentially differs from that
in the Ni-base alloys. It is seen from the results of energy
dispersive X-ray (EDX) microanalysis in a transmission
electron microscope (TEM), Table I, that typical c -solid
solution streng theners, W and Mo, have higher concen-
trations in c 9 th an in c , and Re has s igni ﬁ cant concentration
in c 9 as well, which is different fro m the element partition-
ing in Ni-base superalloys, see e.g., Refs. 19 and 2 0. Results
from previous investigation s
7,10
of the Co-ba se alloy
VIAM-M1, which has s imilar compo sition bu t without Ti
and Cr, are used in this work for comparison of the
c 9 -morphologies. Its composition is given at the bottom
of Table I.
Single crystals of the experimental Co-alloy were
solidi ﬁ ed by the Bridgman method with liquid metal
cooling. The crystal orientation was set by a seed placed
in the crucible bottom, giving [001] oriented crystal
growth. The single-crystal ingots were cylinders of
diameter 15 mm and length 185 mm. The deviation of
the cylinders ’ axes from [001] did not exceed a few
degrees. After casting, the ingots were subjected to
homogenization (1300 °C/15 h), followed by aging
(700 °C/48 h). After heat treatment, the material con-
sisted of a continuous c / c 9 -microstructure and about 10
area% rest eutectics.
Magnitudes of the “ constrained ” c / c 9 -lattice mis ﬁ t,
d c ¼ a c 0  a c

= 0 : 5 a c 0 þ a c

; ð 1 Þ
where a
c
and a
c 9
are the spacings of the mutually
constrained c - and c 9 -lattices, which were measured by
X-ray diffraction (XRD) using 222 re ﬂ ections of Fe K
a
and 004 re ﬂ ections of the Cu K
a
radiation. The area
irradiated by the X-ray beam was about 8 mm
2
, which
covers about 400 dendritic cells, hence the measured
mis ﬁ t value represents the average over the dendritic
structure. It was found that d
c
 1 0.27% (average of
0.24 and 0.30%) in the alloy VIAM-M1. In the alloy
VIAM-M2, d
c
 0% (average of  0.08% and 1 0.07%).
From the heat-treated ingots of VIAM-M2, cylindrical
creep specimens with a gauge diameter of 5 mm and
a gauge length of 25 mm were machined. They were
tested in a lever arm creep machine under a tensile stress
of 196 MPa at 800 °C in air. Two creep tests were
performed, interrupted after 500 h and 550 h, respec-
tively, at creep strains of about 1%.
The deformed alloy was investigated by scanning
electron microscopy (SEM) and TEM under the aspects
of stability of the c / c 9 -microstructure, phase equilibrium,
and deformation micromechanisms. The SEM used is
a Zeiss LEO GEMINI 1530 VP (Carl Zeiss Meditec,
Oberkochen, Germany) equipped with a Bruker x-Flash
TABLE I. Compositions of Co-base alloys VIAM-M2
this work
and
VIAM-M1
11,12
measured by EDX in SEM, and composition of phases
c , c 9 and b in alloy VIAM-M2 measured by EDX in TEM
this work
.
Al Ti Cr Co Ni Mo Ta W Re
VIAM-M2
Alloy (wt%) 6.9 2.1 7.7 53.5 15.9 1.8 2.7 7.9 1.5
Alloy (at.%) 14.9 2.6 8.7 53.1 15.9 1.1 0.9 2.5 0.5
c -phase (at.%) 9.2 0.9 14.8 58.5 12.2 1.2 0.3 2.0 0.8
c 9 -phase (at.%) 13.2 2.6 7.9 53.6 16.5 1.4 1.0 3.3 0.4
c
c 9
/ c
c
1.4 2.9 0.5 0.9 1.4 1.1 4.0 1.6 0.5
b -phase (at.%) 42.9 3.0 2.5 33.5 16.9 0.2 0.3 0.5  0
VIAM-M1
Alloy (wt%) 6.3 ... ... 67.5 15.9 1.0 2.3 5.9 1.0
Alloy (at.%) 13.6 ... ... 66.9 15.9 0.6 0.7 1.9 0.3
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4467
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5030 detector for EDX mic roanalysis and an e

Flash HR
detector for electron backscatter diffraction (EBSD) w ith the
software Esprit and CrystAlign (Bruker, Berlin, Germany).
Specimens for SEM were cut parallel to the longitudinal
planes (100), mechanically gr in ded, and t hen polished . The
ﬁ n al step of preparation was polishing with an emulsion
contai ning 0.2 5 l m silica powder, necessary to remove t he
mechanically deformed surface layer, which would distort
the EBSD pattern. The TEM used is a Philips CM30
(Philips Elec tron Optics, Eindhoven, the Netherla nds) with
an EDX microanalysis system NSS 2.3 from Thermo
(Thermo Electron Corporation, Madison, Wisconsin). The
TEM foils were cut perpendicular to the spe cimen axis
[001] and then electrolytically t h i n n e di na ne l e c t r o p o l i s h e r ,
Struers Tenupol-3 (Struers, C openhagen, Denmark), with
an electrolyte of 10 vol% perch loric acid, 83 vol% ethanol,
and 7 vol% glycerine, at  30 °C. Dislocat ions were
observed in a two-beam bright ﬁ eld diffraction contrast,
and their Burgers vectors b were determined by th e contrast
extinction rule gb 5 0, where g is the reciprocal l attice
vector positioned such that it meets the Bragg condition.
The glide plane was determined by orienting i t in the edge-
on position, so the dislocation g liding on it appears st raight.
III. RESULTS
Figure 1 shows creep curves of the alloy VIAM-M2
obtained in two tensile creep tests at 800 °C, 196 MPa,
interrupted at a creep strain of about 1%. It exposes that
in this strain range, the Co-base alloy deforms with nearly
a constant strain rate of about 5.5  10
 9
s
 1
5 2 
10
 3
%/h. No transient creep is observed for the ﬁ rst
specimen (dashed line); for the second (solid line), one
can discern a very small transient creep of about 0.08%.
Figure 2 shows the microstructures of VIAM-M2 and,
for comparison, VIAM-M1. VIAM-M2 is depicted in the
heat treated state [Fig. 2(a)] and after creep [Figs. 2(b)
and 2(c)], VIAM-M2 after heat treatment [Fig. 2(d)] and
degradation test (stress-free annealing) [Fig. 2(e)]. The
temperature of the degradation test was 800 °C, the same
as under the creep test on VIAM-M2.
Comparing the heat treated microstructures of alloys
VIAM-M1 with nonzero mis ﬁ t [Fig. 2(d)] and VIAM-M2
with zero mis ﬁ t [Fig. 2(a)], it is seen that in the ﬁ rst case,
the c 9 -precipitates have cuboidal morphology, while in
the second case, the c 9 -morphology is polyhedral. In both
cases, the c 9 -precipitates are very ﬁ ne, and their size is
about 100 nm. The c -channels separating the polyhedral
c 9 -precipitates in the alloy VIAM-M2 are very narrow,
and their width is about 10 nm.
Figure 2(e) shows the c / c 9 -microstructure of the alloy
VIAM-M1 after a degradation test (125 h/800 °C), during
which it has become rafted even without stress applied. The
raft thickness is ap proximately equal to the in itial c 9 -size.
The stable plate-like c 9 -morpholo gy nearly does not change
during further ann ealing at 800 °C up to 1000 h.
10
The structural behavior of VIAM-M2 is t otally different .
The initial polyhedral c 9 -morphology in this alloy is
preserved during c reep in s pite of signi ﬁ cant co arsening,
see Fig. 2(b). The c 9 -precipitates coarsen almost isotropi-
cally, and their size after creep scatters around 0.9 l m
(equivalent diameter), which means that the c 9 is coarsened
by almost 10 times! It is remarkable that in local areas near
the eutectic inclusions, the c 9 -morph ology is clearly cuboi-
dal [see Fig. 2(c), left], which is typical for superalloys with
nonzero mis ﬁ t, but with increasing the distance from the
eutectic, the c 9 -precipitates become large r and irregular like
in most volume of the material. Obviously, such a micro-
structural inhomogeneity results from gradients in the
chemical composition in the vic inity of the eutectic, which
is a consequence of den dritic segre gation.
In the alloy VIAM-M2, about 2 – 3 area% of the needle-
shaped phase precipitated in the dendritic arms during
creep, see Fig. 3(a). Precipitates of such a morphology
were observed in the alloy VIAM-M1 after short tensile
tests at 1000 °C,
7
as well as after 1000 h annealing at
800 °C.
10
In Ref. 7, these precipitates were identi ﬁ ed
in TEM as the b -phase with a crystal structure B2
(CsCl type). In the present work, the needle-shaped
precipitates in VIAM-M2 were investigated by EDX in
TEM and EBSD in SEM. The TEM investigations
con ﬁ rmed that the observed precipitates are b -phase
consisting mostly of Al, Co, and Ni and have a stoichi-
ometry close to Al
3
(Co
2
Ni), see Table I.
The relative orientations of the c / c 9 -lattice and the
b -lattices were investigated by EBSD in SEM. Diffrac-
tion patterns were acquired from an area of the specimen,
where patterns from several b -precipitates could be
obtained in addition to those of the single-crystal
c / c 9 -matrix. Figures 3(b) – 3(d) display the pole ﬁ gures
for the standard projections {100}
b
, {111}
b
, and {110}
b
,
which are preferably used as proof for the corresponding
orientation relationship between fcc and body-centred
cubic (bcc) crystal structures. The shown pole distribu-
tions are very close to the orientation-relationship model
proposed by Kurdjumov and Sachs
21
: {111}
c / c 9
k {110}
b
,
FIG. 1. Creep curves of two [001]-oriented single crystals of the
Co-base alloy VIAM-M2 tested at 800 °C/196 MPa.
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4468
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h 110 i
c / c 9
kh 111 i
b
. It assumes that close-packed lattice
planes as well as directions are parallel to each other.
The reference axes x , y , and z are parallel to h 100 i
c / c 9
.
Detailed description of the characterization of the fcc/bcc
orientation relationship by EBSD is given in Ref. 22.
Figure 3 gives the orientation relationship between the
c / c 9 - and b -lattices but not the growth direction of the
needle-shaped b -precipitates. Therefore, identi ﬁ cation of
the axis direction of the b -needles was performed in the
TEM, see Fig. 4. To determine the crystallographic
orientation of the axes of the b -needles, they were
oriented in the upright position (parallel to the electron
beam) by tilting the TEM foil. This orientation is
recognized by the sharp contour of the needle because
the precipitate does not overlap with the matrix. In
Fig. 4(b), needle I is in an upright position and needles
II and III are not. Figure 4(c) shows the corresponding
diffraction pattern of needle I. The 6-fold symmetry of
the spot positions in the diffraction pattern indicates that
the beam is almost parallel to the h 111 i zone axis, which
FIG. 2. c / c 9 -microstructure of Co-base alloys before and after testing. (a, b, c) Alloy VIAM-M2. (a) After heat treatment, TEM. (b) After creep
test (550 h/800 °C/196 MPa), longitudinal cut (100), and SEM. (c) After creep test, cross section (001), and composed TEM image. (d, e) Alloy
VIAM-M1, after heat treatment and after a 125 h degradation test at 800 °C, TEM, quoted from Ref. 10.
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4469
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means that the needle axis coincides with the h 111 i axis
of the b -lattice. The same result was obtained for several
analyzed b -needles. Detailed EBSD analysis showed that
the [111] needle axis of a b -precipitate is parallel to one
of the six h 110 i directions of the c / c 9 -matrix.
Figure 5 shows TEM micrographs of typical disloca-
tion con ﬁ gurations observed in the alloy VIAM-M2 crept
for 550 h at 800 °C and 196 MPa. It is seen from Fig. 5
that dislocations propagate simultaneously in two phases,
c and c 9 . These images were used for identi ﬁ cation of
Burgers vectors b , line vectors u , and glide planes; the
results are presented in Table II. Applying the invisibility
criterion gb 5 0, eight dislocations in Fig. 5
were analyzed. In all cases, Burgers vectors b were found
to be parallel to crystallographic directions h 011 i .I n
Fig. 5(e), the glide plane of dislocations 2 and 3 are in the
edge-on position; it is found to be 
11 
1 ðÞ . The approxi-
mate line vector along the length of dislocations 2 and 3
is 
101 ½ . In Fig. 5(f), a part of dislocation 1 resides in
a glide plane which is in the edge-on position; it is found
to be 
1 
1 
1 ðÞ , and the line direction is 
110 ½ . The glide
planes determined by orienting in the edge-on position
are in accordance with the plane normal calculated by
b  u . From this, we concluded that dislocations glide on
the 01 
1
hi
111
fg
glide system, which is typical for fcc
crystals. The dislocations are mostly of mixed type, with
60° deviation between b and u . Observation of large
areas of TEM foils revealed neither antiphase boundaries
(APBs) nor stacking faults, which could indicate c 9
cutting by partial dislocations. Therefore, it was con-
cluded that the dislocations gliding in c 9 are super-
disloactions a h 011 i . A short matrix segment connecting
two c 9 dislocations over a narrow c -channel is assumed
to be a couple of a /2 h 011 i dislocations, where splitting is
blocked by a strong Orowan force. Entering of the
c -channels by tongue-shaped dislocation loops was
observed very rarely, only in local speci ﬁ c areas with
the coarse c / c 9 -microstructure.
IV. DISCUSSION
It was found that the structural and deformation
behavior of the Co-base alloy VIAM-M2 with zero mis ﬁ t
signi ﬁ cantly differs from that of Co-base and Ni-base
FIG. 3. (a) BSE image of b -needles precipitated in the c / c 9 -microstructure of the Co-base VIAM-M2 alloy after 500 h creep at 800 °C/196 MPa.
(b, c, d) {100}
b
, {111}
b
, and {110}
b
pole ﬁ gures con ﬁ rming the Kurdjumov – Sachs orientation relationship {111}
c / c 9
k {110}
b
, h 110 i
c / c 9
kh 111 i
b
,
EBSD.
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4470
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alloys with nonzero mis ﬁ t. After heat treatment, the
c 9 -phase has a polyhedral morphology. The polyhedral
c 9 -particles coarsen isotropically during creep without
pronounced rafting. The creep curves are linear (constant
strain rate) without essential transient creep. The micro-
mechanism of creep is also speci ﬁ c: Dislocation loops do
not primarily spread along the matrix channels but just
glide simultaneously in both c and c 9 . Although this
structural and deformation behavior is quite unusual, it
ﬁ ts well with the parameters of the c / c 9 -microstructure,
namely zero mis ﬁ t and small channel width. This will be
shown in the discussion of individual ﬁ ndings in the
following four paragraphs.
A. Polyhedral shape after heat treatment
It is known that the cuboidal c 9 -shape with {001} faces
is stabilized by mis ﬁ t stresses. This is because the energy
of mis ﬁ t stresses is minimal when the highly stressed
c -channels have {001} orientation. This, in turn, is
because the lattice stiffness is minimal along h 001 i
directions lying in {001} planes. If the c / c 9 -mis ﬁ ti s
equal to zero, the c 9 -shape is determined by the interface
surface tension instead, which should result in a spherical
shape. But packing of the equal c 9 -spheres is possible
only up to 74% volume fraction or even less if the
c -channels between the c 9 -precipitates still exist.
The c 9 -fraction estimated in the alloy VIAM-M2 from
its composition is close to 80 vol%, hence, it exceeds the
upper critical value. Therefore during c 9 -growth, the
approaching surfaces of neighboring c 9 -precipitates be-
come ﬂ at which leads to a polyhedral shape.
B. Coarsening withou t rafting
It is shown in many publications, e.g., in Refs. 23 and
24 that mis ﬁ t stresses are needed for rafting to take place
during creep. When superalloys with a negative mis ﬁ t are
creep deformed under uniaxial stress, a superposition of
the applied and mis ﬁ t stresses acts on the material. When
the N-channels (i.e., channels oriented normal to the load
axis) undergo plastic deformation, this results in a re-
laxation of the mis ﬁ t stresses in the N-channels, which in
turn results in the formation of an anisotropic pressure
ﬁ eld around the c 9 -precipitates . This anisotropic pres-
sure ﬁ eld activates a cross diffusion of alloying
elements, resulting in an incr ease of the width of the
relaxed N-channels and narrowing of the stressed
P-channels parallel to the load axis. This leads to the
change of c 9 -s ha pe a nd ﬁ nally in coalescence of the
neighboring c 9 -precipitates, i. e ., rafts perpendicular to
the load direction. If the c / c 9 -mis ﬁ ti sz e r o ,t h e n
no mis ﬁ t stresses support the applied stress in the
N-channels; therefore, no prefer able plasti ﬁ cation of
FIG. 4. b -precipitates in the Co-base alloy VIAM-M2 after 550 h creep at 800 °C, 196 MPa (TEM micrographs). (a) Needle lying in the foil ’ s
plane (composite image). (b) Needles passing through the foil, I is in an upright position. (c) Diffraction pattern of needle I.
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4471
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the N-channels takes place, and ﬁ nally no drivin g force
for cross diffusion and rafting evolves.
C. Cuboidal c 9 -shap e near the eutec tics
It was shown in Refs. 16 and 25 that segregation of
alloying elements in Ni-base alloys during dendritic growth
results in inhomogeneity of the c / c 9 -mis ﬁ t with in a dend ritic
cell. The degree of dendritic segregation depends on the
alloy composition and somewhat differs for Ni- and C o-base
alloys. However, the segregation char acter is similar for
alloys of these two classes, e.g., in Ni-as well as in Co-base
alloys, W segregates into the de ndritic arms, while Ta
segregates into the interdendr itic regions, see e .g., Ref. 9.
In Ref. 26, quaternary Co – Al – W – Ta alloys were character-
ized; in particular, the effect of alloy composition on spacing
of the c -a n d c 9 -lattices, a
c
and a
c 9
, was investigated. F itting
of the results gave the following linear equations:
a c 0 nm ½ ¼ 0 : 35976
þ 5 : 1 c Al þ 2 : 2 c W þ 26 : 4 c Ta
ðÞ  10  5 ;
ð 2 Þ
a c nm ½ ¼ 0 : 35694
þ 3 : 3 c Al þ 17 : 1 c W  2 : 9 c Ta
ðÞ  10  5 ; ð 3 Þ
where c
i
is the concentration of the i th element in at.%.
It ’ s seen from Eq. (2) that W signi ﬁ cantly increases the
c -spacing a
c
(preconcentration factor 17.1  10
 5
nm/at. %),
while from Eq. (3) it can be followed that Ta signi ﬁ cantly
increases the c 9 -spacing a
c 9
(preconcentration factor
26.4  10
 5
nm/at.%). Taking into account the de ﬁ nition
of the c / c 9 -mis ﬁ t, Eq. (1), one can see that Ta
FIG. 5. Dislocation con ﬁ gurations after 550 h creep at 800 °C,
196 MPa used for identi ﬁ cation of Burgers vectors, line vectors, and
glide planes (TEM micrographs). Images (a-h) present the same area
but under different imaging conditions ( k
0
and g ) given in Table II.
Note that here, the lattice direction parallel to the stress axis is called
[010]. Arrows give the projections of lattice vectors into the image
plane. In (d), the orientation of the lattice ’ s unit cell is indicated by
a cube.
TABLE II. Identi ﬁ cation of Burgers vectors from dislocation contrasts
in the TEM. ● stands for visible and s for extinct dislocation contrast.
Note that here, the lattice direction parallel to the stress axis is called
[010]. Line vectors are only approximate.
Image no.
Imaging
conditions Contrast of dislocation
k
0
g 12 3
5a [010] [002] ● ss
5b [010] 
200
½ s ●●
5c [010] 20 
2 ½ ●● ●
5d [010] [202] ●● ●
5e [121] 1 
11 ½ ● ss
5f 
12 
1 ½ 
1 
1 
1 ½ s ●●
5g 
121 ½ [202] ●● ●
5h 12 
1 ½ 
20 
2 ½ ●● ●
6 b parallel to 01 
1 ½ [110] [110]
Line direction u 
110 ½ 
110 ½ 
110 ½
Angle ( b , u ) 60° 60° 60°
Glide plane (111) 
11 
1 ðÞ 
11 
1 ðÞ
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4472
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signi ﬁ cantly increases the mis ﬁ t value d
c
, while W
signi ﬁ cantly decreases d
c
. Thus, the average mis ﬁ ti n
the alloy VIAM-M2 is close to zero, but in the interden-
dritic region, it can increase due to higher c
Ta
, as well as
lower c
W
. This local increase of the mis ﬁ t is obviously
the reason why the c / c 9 -microstructure is cuboidal and
not polyhedral in the interdendritic regions of the VIAM-
M2 alloy, that is, close to the eutectic, see Fig. 2(c).
D. Constant-rate creep curves and dislocation
micromecha nism of deformation
Under the testing condition s used, the Co-base alloy
VIAM-M2 deforms by the glide of dislocation loops on the
h 011 i {111} glide system, as is typical for fcc crysta ls. It is
remarkable that the dislocation l oops glide simultaneously
in both c and c 9 , keeping a strai ght shap e when they c ross
the c -channels. This is different from the dislocation
mechanism observed in Ni- base supe ralloys w ith noze ro
mis ﬁ t, where creep deformat ion starts by gliding of the
dislocation loops in the matrix channels, see, e.g., Refs. 27 –
29. This primary matrix glide results in a fast transient
creep, which then slows down due to a back s tress of
deposited interfacial dislocat ions. In VIAM-M2, dislocation
loops glidin g exclusively in th e matrix channels and
dislocations deposited in the c / c 9 -interface by the matrix
glide were very rarely observed. This suggests that the
primary matrix glide was not genera lly activated in the
investigated Co-base alloy, and this e xplains why transient
creep is not observed. There are two obvious reasons for
this: Firs t , in an alloy wit h zero mis ﬁ t, there is no m is ﬁ t
s t r e s s supporting the entering of dislocation loops into
the matrix channels. Second, without rafting rapidly
widening the matrix channels, the matrix glide can be
blocked by a strong Orowan back stress for an
extended time, in our case for 500 h. However, at
longer creep times, isotropic coarsening will widen the
c -channels, permitting the matrix glid e. Such a struc-
tural behav ior (without rafti n g )a sw e l la sad e f o r m a t i o n
behavior (without transient creep) was observed in an
experimental Ni-base alloy UM-F22 which has
c / c 9 -mis ﬁ t close to zero.
30
During testing at 950 °C
and 290 MPa, this alloy showed no rafting and slow
creep for 100 h. But after this time, the
c / c 9 -microstructure became signi ﬁ cantly coarsened,
and the creep rate rapidly accelerate s. Deformation
without the primary m atrix glide was reported in Re f.
31 for a high-temperature low-stress creep (1100 °C
and 137 MPa) of a superalloy TMS-75 with a small
lattice mis ﬁ t. Due to the small value of c / c 9 -lattice
mis ﬁ t, the dislocations coul d not glide in the matri x
channels but mainly climbed around the c 9 -cuboids. In
our case, such a climbing is not activated due to a much
slower diffusion at 800 °C.
c 9 cutting by h 011 i superdislocations is a classical
mechanism ﬁ rst reported by Kear and Wilsdorf for the
L1
2
structure of Cu
3
Au
32
and then reported for Ni-base
superalloys many times, see e.g., Ref. 33. However, this
me ch an is m i de nt i ﬁ ed in our wo rk diff er s fr om tha t
re por te d re ce ntl y fo r the c reep of Co -b ase al loys a t
900 °C.
13 – 15
Th e auth ors fo und c 9 cu ttin g by a /2 h 01 1 i
an d a /2 h 112 i dis loca tion s, ac com pan ied b y the f orma tio n
of A PB s an d sta ckin g fau lts . He re , on e sh ould n otic e that
a str ess le vel ap pl ie d in ou r case, 19 6 MPa , is sig ni ﬁ can tly
lo wer th an str ess le vels ap pl ied in Ref s. 13 – 15, be twee n
275 an d 345 M Pa. It is we ll kno wn th at hig h stres se s are
nee ded to pre ss th e par tial d islo cati ons ins ide a s uper latt ice.
Mor eov er, at 800 °C, the te mper atu re of ou r cree p tes ts,
th e Co -bas e allo ys s how the m axim um yie ld str es s,
3
wh ich
in dic at es a hi gh she ar str engt h of th e c 9 -pha se.
V. CONCLUSION S
( 1) An experimental Co-base alloy with zero
c / c 9 -lattice mis ﬁ t showed unusual structur al and de-
formation behavior, namely: After heat treatment, the
c 9 -phase in this alloy has a poly hedral morphology.
During creep at 800 °C and 196 MPa, the polyhedral
c 9 -precipitates coarsen is otropically without pro-
nounced rafting. Tests at 800 °C and 196 MPa
performed up to about 1% creep strain give creep
curves with a constant strain rate and without transient
creep. Creep micromecha nism is a glide of dislocations
simultaneously in the two phases, c and c 9 , w ithout
preliminary spreading of dislo cation loops along the
matrix channels. All these features can be well
explained by the zero c / c 9 -lat tice mis ﬁ ta n d ﬁ ne
c / c 9 -microstructure.
(2) It is reasonable to ask: does an alloy with zero
mis ﬁ t have advantages compared with an alloy with
nonzero mis ﬁ t? One obvious advantage is the absence of
a fast transient creep which is caused by two factors: (i)
Zero mis ﬁ t stresses and (ii) thinner channels during
isotropic coarsening (without rafting). However, these
factors grant a positive contribution only at the initial
stage of the creep.
In superalloys with nonzero mis ﬁ t, the mis ﬁ t stresses
initially promote dislocation glide, but relax during
primary creep. Subsequently, even back stresses are
induced, which retard dislocation glide in the matrix.
Rafting in these alloys rapidly widens the matrix chan-
nels, but then, the channel widening slows down because
the rafts are stable. In zero mis ﬁ t alloys, however, the
polyhedral c 9 -precipitates coarsen continuously. There-
fore, one should expect that after a certain time, the
channel width in their polyhedral microstructure can be
larger than in rafted microstructures, hence their relative
advantage would be lost.
(3) During high-temperature exposure, the phase trans-
formation c 9 ! b occurs in Co-base alloys. It was found
that the b -needles grow in the h 111 i directions of the B2
J. Midtlyng et al.: Creep behavior of a c 9 -strengthened Co-base alloy with zero c / c 9 -lattice misfit at 800 ˚C, 196 MPa
J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4473
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lattice, which correspond to the h 110 i directions of the
c / c 9 -matrix. The b -precipitation after 550 h creep at
800 °C/196 MPa was not found to be strong, 2 – 3 area
% in the dendrite arms.
ACKNOWLE DGMENTS
The authors are grateful to the German Research
Foundation (DFG), project EP 136/2-1, NO 307/5-1
and the Russian Foundation of Basic Research (RFBR),
project 13-08-91330- ННИО _ а for funding this work.
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J. Mater. Res., Vol. 32, No. 24, Dec 28, 2017 4474
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