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International Journal of Fatigue
journal homepage: www.elsevier.com/locate/ijfatigue
On the mechanistic difference between in-phase and out-of-phase thermo-
mechanical fatigue crack growth
V. Norman
a,⁎
, S. Stekovic
a
, J. Jones
b
, M. Whittaker
b
, B. Grant
c
a
Division of Engineering Materials, Department of Management and Engineering, Linköping University, SE-58183 Linköping, Sweden
b
Institute of Structural Materials, Swansea University, Swansea SA1 8EN, UK
c
Rolls-Royce plc, Derby DE24 8BJ, UK
ARTICLE INFO
Keywords:
Aerospace
Superalloys
Thermomechanical fatigue
Crack growth rate
Crack opening
ABSTRACT
The crack driving mechanisms in a coarse grained nickel-base superalloy RR1000 when subjected to in- and out
of phase thermo mechanical fatigue are investigated. It is found that the difference in fatigue crack growth rate
between these two load conditions is accounted for by the different mechanical conditions at the crack tip
region, rather than oxidation effects. This is based on digital image correlation and finite element analyses of the
mechanical strain field at the crack tip, which demonstrate that in phase leads to larger crack tip deformation
and crack opening. Notably, it is demonstrated that in- and out of phase crack growth rates coincide when
correlated to the crack tip opening displacement.
1. Introduction
In recent years, there has been an increased awareness of the en-
vironmental impact of air travel. Therefore, it is of vital strategic im-
portance to the aviation industry to reduce aero engine emissions,
which is driven by the EU’s Aviation Vision 2020 along with industrial
competition [1].
A significant portion of these improvements are expected to come
from new engine designs in term of gas turbine efficiency. Such an
increase in efficiency of the gas turbine is usually achieved either by
weight reductions or by increasing the combustion temperature as a
result of fuel being burnt at temperatures approaching the stoichio-
metric value [2]. In either case, the material choice is of critical im-
portance. Effectively, the modern criteria for selecting materials include
requirements on high temperature fatigue and creep capabilities, as
well as requirements on suitable environmental and corrosion resistant
properties, which in the present context typically results in the em-
ployment of nickel base superalloys. The motivation of these strict
criteria is that the gas turbine operation cycle imposes harsh alternating
mechanical and thermal loads on the material during start up, take off,
descent and shut down. Such thermo mechanical conditions have the
potential to cause local stresses to peak at temperatures far below the
flight cycle maximum, resulting in the nucleation and propagation of
cracks; a phenomenon known as thermo mechanical fatigue (TMF).
In essence, TMF is a complex failure mechanism, caused by
combined thermal and mechanical load cycles [3], but is not a novel
phenomenon. Rather, it is well acknowledged to be the primary life
limiting aspect for many engineering components exposed to elevated
temperatures, such as parts in the combustion chamber, along with
turbine blades and discs [4,5]. More precisely, the TMF process can be
divided into an initiation and propagation stage, where this investiga-
tion addresses the latter. Only a very limited number of investigations
have been conducted involving TMF crack growth experiments on
nickel base superalloys, however, it has been established that the crack
growth rate is highly dependent on the phase angle between the
thermal and mechanical cycle, e.g. in (IP) and out of phase (OP) cycling
[6–8]. This introduces new perspectives regarding the general crack
growth mechanism in fatigue at elevated temperatures, which have
been extensively studied in the past based on isothermal fatigue and
dwell fatigue experiments [9–27]. In particular, based on crack growth
experiments in air and vacuum [9–17], the general conclusion is that
oxygen plays a significant role, possibly through one of a number of
proposed mechanisms [16]. On top of this, some authors have argued
that inelastic creep deformation and stress relaxation at the crack tip
may influence the growth rate [7,11,12,14,15,19,20]. However, it is
not yet known whether these mechanisms are able to account for the
differences seen between IP and OP TMF crack growth experiments.
In view of the above investigations, three different kinds of me-
chanisms which potentially may account for the difference in OP and
IP, are identified. Firstly, (i) it is suggested that crack closure effects
https://doi.org/10.1016/j.ijfatigue.2020.105528
Received 1 November 2019; Received in revised form 31 January 2020; Accepted 1 February 2020
⁎
Corresponding author.
E-mail address: viktor.norman@liu.se (V. Norman).
International Journal of Fatigue 135 (2020) 105528
Available online 08 February 2020
0142-1123/ © 2020 Published by Elsevier Ltd.
T
may influence the crack growth rate depending on the thermo me-
chanical phase angle. For instance, recent studies [8,28–30] have de-
monstrated that variations in crack growth rate caused by altering the
load ratio and temperature cycle are accounted for by compensating the
stress intensity factor range or the cyclic J integral with respect to crack
closure. Secondly, (ii) it is also expected that the different phase angles
may induce different stress strain states in the crack tip, which may
both explicitly and implicitly affect crack growth. For instance, dif-
ferent amount of inelastic crack tip deformation may occur depending
on whether the cycle is OP or IP. Conversely, implicit effects such as the
relation between the stress strain state at the crack tip and the diffusion
of oxygen are also possible [9,10,16,31]. Accordingly, as the third and
final category, (iii) it is also likely that aspects related to the material
structure may account for the effect of phase angle, including the
mentioned environment material interaction and potential crack tip
phase transformations [11,32]. In other words, any potential material
related weakening or toughening at the crack tip, dependent or in-
dependent of the atmosphere.
Even though fatigue crack growth mechanisms in nickel base su-
peralloys under isothermal conditions have been extensively in-
vestigated in the past, there is at present a very limited understanding
of the governing mechanisms under thermo mechanical load condi-
tions, which is necessary for the efficient development of next genera-
tion gas turbine materials. Accordingly, the objective of this in-
vestigation is to render a better understanding of the mechanisms
responsible for causing the difference in crack growth rate between IP
and OP TMF load conditions. This is done by performing IP and OP TMF
crack growth tests on a coarse grained nickel base superalloy RR1000,
for which the potential mechanisms, including crack closure, crack tip
deformation and oxidation effects, are investigated by in situ digital
image correlation (DIC) and metallographic studies of interrupted tests.
In this way, the mechanisms are identified and their contribution as-
sessed by comparison of the IP and OP case. This has given new insights
regarding general aspect of crack growth at elevated temperatures re-
levant for both isothermal and thermo mechanical load conditions.
2. Experimental and computational methods
2.1. Materials
RR1000 is a superalloy developed by Rolls Royce plc and is mainly
used for discs in the rotative aero engine hot sections because of its
excellent high temperature mechanical properties. The fine grained
RR1000 has at least 25 °C increase in temperature capability over 720Li
and an equivalent crack growth behaviour to coarse grained Waspaloy
[33]. However, the need for increased engine efficiency and lower
emissions has been driving towards higher pressure ratios and higher
operation temperatures. Therefore, a coarse grained RR1000 has been
developed to improve creep performance and fatigue crack growth re-
sistance of these components [34].
Coarse grained RR1000 is processed by a powder metallurgy route
and strengthened by
N
iA
l
3type gamma prime precipitates. The nominal
chemical composition in wt% and details about the processing as well
as the heat treatment of the alloy are given in reference [14]. The mi-
crostructure typically consists of grain sizes of ASTM 7 3 (32–125 μ
m
).
Apart from the gamma matrix and main strengthening secondary and
tertiary gamma prime phases, the material also contains a dispersed sub
micron phases such as MC carbides and
M
B
3
2
borides. Both gamma
prime precipitates are distributed intergranularly with particle sizes
ranging from 1 μ
m
to 5 nm depending on the heat treatment and
cooling rates [35,36].
2.2. Measurement of fatigue crack growth rates under thermo mechanical
fatigue load conditions
The purpose with this test procedure is to assess the fatigue crack
growth rate when subjected to thermo mechanical cycling, in contrast
to standard fatigue crack growth tests during which the temperature is
constant. As similar to conventional thermo mechanical fatigue (TMF)
testing [3], the temperature varies with the same periodicity as the
applied mechanical load with a given phase angle such as for instance 0
or 180 degrees, often denoted as in phase (IP) and out of phase (OP),
respectively. As a consequence of the varying temperature, a distinction
between the different causes of deformation in the tests specimen must
be made, typically as
Fig. 1. (a) The single edge notched and (b) corner crack specimen geometry used for the TMF crack growth experiments.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
2
=+εt ε t ε t() () (
)
ethmech
(1)
where
ε
e
is the actual strain measured by an extensometer,
ε
th
is the
thermal strain, i.e. the thermal expansion which varies with the tem-
perature, and εmech is the mechanical strain whose origin is solely the
applied force.
2.2.1. Specimen geometry
For the aforementioned purpose, the specimens used were notched
in order to initiate a starting crack following a pre cracking procedure
explained below. The specimen geometry is displayed in Fig. 1a. The
grip section of the specimen was cylindrical with a diameter of 12 mm
while the middle section had an approximately rectangular cross sec-
tion with a thickness of 3 mm and a width of 12 mm. Nominally, the
notch had a radius of 1 mm and a notch depth of 3.0 mm. The speci-
mens were manufactured through turning and wire electrical discharge
machining, without application of any additional surface finishing
process.
2.2.2. Pre cracking procedure
A pre cracking procedure was conducted to initiate and propagate a
starting crack to a reasonable length. Due to lack of experience, the pre
cracking procedure was varied for the first specimens until an optimised
procedure regarding minimal duration and applied stress, was estab-
lished. For this reason, the procedure was not the same for most of the
specimens. The load parameters and notch geometry employed for
crack initiation are displayed in Table 1. Subsequent propagation to
reach a reasonable starting crack length was done in a stepwise manner
reducing the maximum stress with steps of 30 MPa in order to propa-
gate through any eventual plastic zone caused by the previous max-
imum stress value. Regarding the test initiated at
σ
max
= 210 MPa and
R=−1(
S
3
and
S5
in Table 1), these were subsequently propagated at
240 MPa, and then again at 210 MPa with R = 0 for the same reason
prior to the actual test. At the end of the pre cracking procedure for
each specimen, the crack had a length of about one millimetre. It should
also be noted that one of the tests,
S
4
was initiated and propagated
using the same temperature cycle as the actual test with the purpose of
investigating the effect of a TMF initiated starting crack as similar to
what is expected in a real component. Furthermore, the second test
S
2
was carried out with an alternative notch appearance, which however,
did not decrease the duration of the pre crack procedure.
2.2.3. Test procedure
After having reached the desired starting crack length, the crack
growth tests were conducted in both a IP and OP configuration, with a
temperature cycle between 400 and 750 °C. The total cycle time was
70 s divided into 35 s ramp up and down in temperature with a tri-
angular wave shape. The mechanical load consisted of a prescribed
stress with the same periodicity as the above cycle and a load ratio R
equal to zero. A total number of five crack growth test were conducted
as presented in Table 1.
2.2.4. Test set up
All tests and pre crack operations were conducted in an Instron
8801 servo hydraulic test machine equipped with an induction heating
system including a cylindrical copper coil with its centre axis coinciding
with the specimen centre axis. To even out the temperature distribution
and assist cooling, compressed air flow directed towards the specimen
through three nozzles was used, positioned circumferentially with equal
angular spacing and the same distance to the specimen, as well as at the
same vertical position as the notch. All tests were controlled and
monitored using a dedicated TMF software developed by Instron, which
automatically performs a pre test procedure, involving thermal stabi-
lisation, thermal strain measurement and validation. In addition, an
elastic modulus measurement program included in the software was run
prior to the test on each specimen, measuring the elastic modulus at
different constant temperatures using load cycles of
±
20 MPa, hence
safely within the elastic range of the material. The strain was measured
using an Instron extensometer 2632-055 with a gauge length
L
e
of
12.5 mm and the temperature was measured using a N type thermo-
couple spot welded in the centre of the side surface of the specimen,
slightly beneath the expected crack path. After each test, it was verified
that the thermocouple did not interfere with the crack path.
Before the start of the test series, a thermal profiling procedure was
conducted in order to evaluate the temperature distribution on the
specimen. For this purpose, a dummy specimen was tested to which six
different N type thermocouples were attached at six different locations;
three on each side evenly distributed along the axial centre line of the
specimen. By monitoring the temperature at each thermocouple, the
coil and air nozzles were adjusted in order to achieve a temperature
difference less than 10 °C throughout the selected temperature cycle, as
recommended by the TMF standard [3]. Evidently, the chosen config-
uration of the coil and air nozzles was not changed during the sub-
sequent test series.
2.2.5. Crack length measurement method
The crack length as a function of number of cycles was determined
using the compliance method in accordance with previous investiga-
tions involving the same or similar specimen geometry [29,37,38]. The
method is based on the correlation between the crack length and the
compliance or stiffness of the specimen, which is acquired from finite
element simulations of the employed specimen geometry. The method
is outlined in detail in [38] and is only briefly reproduced here.
Afinite element model representing the specimen was set up using
Abaqus CAE version 6.12 with the nominal specimen dimensions
mentioned above, see Fig. 1a. The specimen has two symmetry planes
which can be exploited, therefore only one quarter of the specimen was
included in the model. In addition to the symmetry boundary condi-
tions, uniaxial and monotonic traction was prescribed on the cross
section situated at the grips by restricting the node displacement on this
plane to be uniform with a non zero component only in the tensile
direction using a coupling equation to a dummy node [39]. The pre-
scribed stress ramp going up to 200 MPa was then implemented by
assigning an isolated force, whose magnitude is compensated by the
cross sectional area, to the dummy node. Regarding the material
properties, they were only assigned as elastic with data originating from
a tensile test program conducted on the considered materials by Rolls
Royce plc. In total, the model consisted of roughly 6000 brick elements
employed with reduced integration and an approximate size of 200 μ
m
in the gauge length volume.
Starting from the above described model, a plane crack was added
by suppressing the symmetry condition over the surface at which the
Table 1
Information regarding the conducted crack growth tests and the associated pre
cracking procedure.
Name Test Crack initiation Notch geometry
400–750 °C 25 °C
R=0 20Hz
S1
σ
ma
x
= 250 MPa R = 0
IP
σ
ma
x
= 300 MPa
S2
σ
ma
x
= 210 MPa R = 0
IP
σ
ma
x
= 300 MPa
S3
σ
ma
x
= 210 MPa R = −1
OP
σ
ma
x
= 210 MPa
S4
σ
ma
x
= 210 MPa 400–750 °C
IP R = −1
σ
ma
x
= 240 MPa
S5
σ
ma
x
= 210 MPa R = −1
IP
σ
ma
x
= 210 MPa
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
3
crack extension was anticipated. Accordingly, a set of different models
were set up, each having a plane crack of different extension ranging
from 0.5 to 5 mm measured from the notch root. Around the crack tip,
the mesh was refined to a spider web configuration with decreasing
element size closer to the crack tip reaching the smallest size of 10 μ
m
.
The brick elements of the innermost ring were collapsed into wedge
elements whose crack tip nodes were tied. In order to simulate the
experiments, the stiffness of each such model was then assessed as the
slope of the applied stress at the dummy node and the strain, evaluated
as the node displacement at the point where the extensometer arm is
placed divided by the gauge length. The result of this procedure is
displayed in Fig. 2a, where the stiffness normalised to the zero crack
length stiffness is plotted against the prescribed crack extension, in-
cluding a polynomial fit.
Experimentally, the stiffness was evaluated in every load cycle as
the slope of the stress and mechanical strain curve at the turning point
of maximum stress over an interval of 80–95% of the maximum stress
value. This interval corresponds to a temperature interval of 680 to 730
°C over which the elastic modulus of the material varies negligibly. The
un cracked stiffness used to normalised the experimental stiffness value
was taken from the initial stiffness measurement, performed prior to pre
cracking, see Section 2.2.4, at the average temperature corresponding
to the above stress range. Subsequently, using the polynomial expres-
sion given by the FE modelling, the stiffness was converted into a crack
length and then differentiated to obtain the crack growth rate.
2.2.6. Stress intensity factor assessment method
Regarding crack growth, a conventional procedure is to relate the
crack growth rate to the mode I stress intensity factor, here denoted as
K
I
, which is a parameter dependent on the applied stress and crack
length as
=
K
σYaπa()
Imax (2)
where
σ
max
is the nominal stress applied to the specimen and
Ya(
)
is a
geometrical factor dependent on the crack length a.
The geometrical factor Yof this particular load geometry was as-
sessed using the same set of finite element models as described above.
Based on Eq. (2), the geometrical factor was computed as
̃
̃
=Ya K
σπa
() I
max (3)
where
̃
K
I
is the stress intensity factor of the finite element model
evaluated as the average stress intensity factor along the whole crack
front,
̃
σ
max
is the applied stress of 200 MPa and ais the crack length
implemented in the model as mentioned above, Section 2.2.5.By
varying the crack length over the set of finite element models, a func-
tional dependence of Ywas acquired, see Fig. 2b, which was fitted using
a polynomial expression as similar to the stiffness curve.
As will be seen in this investigation, crack closure has an important
influence on the crack growth behaviour. For this reason, the effective
stress intensity factor is computed and correlated to the crack growth
rate, with the purpose of compensating for crack closure [40]. This
parameter is obtained as
=−
K
Yσ σ πa()
Ieff max op,(4)
where
σ
max
is the nominal stress and
σ
o
p
is the crack opening stress
which can be measured experimentally, see Sections 2.3 and 2.5,Yis
the geometrical factor and athe crack length.
2.2.7. Additional thermo mechanical fatigue crack growth tests of corner
cracked specimens
In order to validate the TMF crack growth experiments, additional
tests performed using a different laboratory set up, but with the same
test parameters, were included in the investigation. For these tests, a
corner cracked specimen design with a rectangular cross section was
employed, see Fig. 1b. The test piece gauge length had a 7 × 7 mm
square cross section and 20 mm length, with a 0.35 mm ± 0.01 mm
single edge notch machined using a diamond edge saw blade at one
corner.
The TMF crack growth testing of this specimen type was undertaken
using two different heating methods. The first set up comprised of an
Instron 100 kN servo hydraulic test frame, utilising a Zwick CUBAS
control system. A Trueheat 10 kW induction heating system was utilised
to deliver rapid heating rates through a 4 mm diameter copper tube non
uniform multi turn longitudinal field helical coil with an approximate
external diameter of 60 mm. Rapid cooling rates were enabled through
forced air cooling using four Meech pneumatic air amplifiers with their
output control through proportional solenoid valves. Temperature
feedback was provided through a 0.2 mm diameter N type thermo-
couple, spot welded at the centre of the 20 mm gauge length on an
opposing face to the starter notch.
The second set up consisted of a Instron 100 kN servo-electric test
frame together with a DARTEC control system. A second generation
radiant lamp furnace (RLF) was designed in collaboration with Severn
Thermal Solutions Ltd to generate rapid heating rates similar to that of
an induction coil. The 12 kW RLF was a standard split body design with
each half containing three horizontally mounted lamps. Three in-
dependently controllable heating zones allowed the accurate tempera-
ture control and profiling, whist built in internal compressed air cooling
delivered the rapid cooling rates required by the complex TMF wave-
forms. Again temperature feedback was provided through a 0.2 mm
diameter N type thermocouple, spot welded at the centre of the 20 mm
gauge length on an opposing face to the starter notch.
A Dirlik control system interfaced together with the set ups de-
scribed above was employed to record crack length against number of
cycles readings through pulsing a 10 A signal and utilising the direct
current potential drop technique. The crack length was converted into
crack growth rate data by using the incremental polynomial method, as
described in the ASTM 647 appendixes, standard test method for
measurement of fatigue crack growth rates [41].
Similar to the single edge notched specimen described above, rig-
orous thermal profiling was undertaken using six 0.2 mm diameter N
Fig. 2. Estimated dependence of the (a) normalised stiffness and (b) the geometrical factor Y on the crack length from the conducted FE simulations. Note that
definition of crack length here does not include the notch depth.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
4
Type thermocouples. These thermocouples were spot weld at the centre
gauge location on each of the four rectangular specimen faces to realise
the radial heating gradient. A further two thermocouples were spot
welded 5 mm above and below a centre thermocouple to generate axial
temperature distributions. Similarly to as described in Section 2.2.4, the
authors employed the stringent temperature limits imposed by the
governing TMF strain control standard [3].
The test parameters employed were the same as for the single edge
notched specimen, namely a triangle waveform with =
R
0over a 400
750 °C temperature range with heating and cooling rates at 10 °C/s.
However, due to the difference in the geometry of the specimens, a peak
stress of 500 MPa were employed in order to have similar value in stress
intensity factors. Regarding pre cracking, it was performed at room
temperature using sinusoidal wave with =
R
0.
1
, starting with propa-
gation at 600 MPa and 4 Hz, then 550 MPa and 2 Hz and finally at
500 MPa and 1 Hz, in order to reach the starting crack length gently
thereby avoiding the introduction of any residual plastic strain ahead of
the crack tip. The same pre cracking procedure was used for all tested
corner cracked specimens.
2.3. Crack closure stress measurement method based on specimen stiffness
Crack closure stress is vaguely defined as the nominal stress at
which the crack is closed and has been demonstrated to influence the
TMF crack growth rate [8,28,29]. The effect of crack closure is easily
demonstrated by looking at the cyclic stress strain data in which the
stiffness value significantly changes depending on whether the crack
faces are in mechanical contact or not. Based on this feature, the crack
closure stress was determined in the present study following a sug-
gested procedure outlined by Palmert et al. [8].
Accordingly, the crack closure stress was assessed in terms of the
ratio between the stiffness of an arbitrary cracked configuration Cand
the stiffness of the un cracked reference configuration
C
0
, where both
stiffness variables depend on the instantaneous temperature and nom-
inal stress. By this definition,
CC/
0
supposedly becomes unity when the
crack is completely closed, since the stiffness is expected to be the same
then as for an un cracked specimen, while less than unity when the
crack is open. For mathematical convenience, the stiffness ratio is
transformed by the following operation
=−
−=
DCC
CC
1/
1/|
σσ
0
0max (5)
where =
CC/|
σσ0max is the value of the stiffness ratio at the instant of
maximum nominal stress. In this way, the newly derived crack closure
factor Dis zero when completely closed, and unity when completely
open. Note that this interpretation presupposes that the crack is com-
pletely open at maximum nominal stress, which however is expected to
be true for all tests conducted in this investigation. The variation in
factor Dwill clearly be influenced by the continuous separation of the
crack faces when loaded, since the amount of partial crack closure re-
lates to the specimen stiffness. Accordingly, the crack faces at the crack
tip will be the last to separate and it is therefore motivated to associate
crack closure event with the nominal stress applied when the factor Dis
very close to unity. For this reason, crack closure stress defined as the
nominal stress at which Dexceeds 0.9 in this investigation, as suggested
by Palmert et al. [8].
Experimentally, crack closure was only assessed for the single edge
notched specimen. The stiffness Cwas determined at twenty stress
values regularly distributed over the loading branch of the hysteresis
loop, as the slope of the stress strain curve over an interval of ±7.5% of
the maximum stress of the cycle. The temperature variation over this
interval is 50 °C, over which the variation in elastic modulus is negli-
gibly small. The reference stiffness
C
0
was taken from the initial stiff-
ness measurement, performed prior to pre cracking, see Section 2.2.4,
at the average temperature corresponding to the above stress range.
2.4. Metallographic analysis of interrupted tests
A metallographic investigation was conducted on the specimens
subjected to the OP and IP condition described above, namely specimen
S
3
and
S5
in Table 1 respectively, which were interrupted at the same
crack length of 4.2 mm. The specimens were then cut in order to re-
trieve the rectangular middle section, which in turn were cut with a
cutting plane perpendicular to the crack face and parallel with the
width dimension. In this way, the crack tips were studied both at the
side surface of the specimen and on a plane in the centre of the notch,
i.e. at half the thickness. The metallographic surfaces were then ground
and polished using a standard program for nickel base superalloys.
The microscope equipment used were an optical microscope and a
scanning electron microscope (SEM). The former was a Nikon Optiphot
optical microscope and the latter a HITACHI SU-70 field emission gun
SEM, equipped with a solid state 4 quadrant backscattered electron
detector, using 8 and 10 kV acceleration voltage and a working distance
of about 9 mm. Furthermore, energy-dispersive X-ray spectroscopy
(EDS) was performed at 20 kV at a working distance of 15 mm.
2.5. Image analysis of the crack tip region
In this investigation, the deformation field ahead of the crack tip in
the single edge notched specimen was measured using digital image
correlation (DIC) in order to acquire a better understanding of how IP
and OP loading affect the mechanical conditions at the crack tip region.
For the same purpose, the displacement field was further used to assess
the crack tip opening displacement.
The images were captured using a Nikon UBS29 QXC F camera
mounted at a lateral viewpoint of the specimen. The images were of a
size of 2592 × 1944 pixels and captured at a magnification of about
40×. The camera was positioned to view the notch from a lateral di-
rection and images were captured every 35 s, i.e. twice every cycle at
the instant of maximum and minimum stress. Only for a few occasions,
images were captured more frequently at a frequency of 1 Hz, namely
when the crack length was about 3 and 4.2 mm. No artificially added
speckle pattern was utilised since the natural surface roughness was
enough to acquire accurate correlation in the DIC analyses. Even so,
surface oxidation was not an problem since the time between correlated
images was less than required to produce significant changes in the
appearance of the specimen surface. Furthermore, the view of the
specimen was constantly illuminated using high power LED spots
comparable to 150 W halogen light sources in order to eliminate the
disturbance from the black body radiation of the specimen. An open
source matlab based DIC code written by Eberl et al. at the John
Hopkins University and distributed by mathworks [42], was used for
the image correlation with a subset pixel size of 31 × 31 pixels.
For all conducted DIC analyses, the reference image was selected as
the image taken at minimum load, i.e. zero applied stress, at the same
load cycle as the analysed image of interest. For this reason, the dis-
placement fields presented in this investigation are always with respect
to the deformation state at zero nominal load, which is not necessarily a
state of zero residual deformations.
Within the DIC software, the obtained displacement field was
smoothed, using a Gaussian distribution of weights with a Gaussian
kernel size of 31 control points and three smoothing passes. The
smoothed displacement field was then subsequently differentiated in
order to obtain the strain field. Furthermore, by assuming homo-
geneous temperature in the region of interest, the mechanical strain
was acquired using Eq. (1),i.e. by subtraction of the thermal strain
measured in the TMF pre test procedure.
The crack tip opening displacement (CTOD) was calculated by post
processing of the DIC analyses similar to the method presented by Vasco
Olmo et al. [43]. More precisely, CTOD was defined as the dis-
continuous jump in the field of the displacement component of the
tensile direction (y-direction), at an x-position along the crack 10 μ
m
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
5
from the crack tip location at the instant of maximum nominal stress.
The discontinuous jump in displacement was assessed by fitting a step
function, namely the error function as
=+ −
f
yaberfcyd() · [·( )]
(6)
where abc,, and dare fitting constants, to the displacement component
profile in the y-direction along y-position y. From the fitted parameters,
CTOD was hence defined as b2, since the total step height of the error
function is 2. This procedure was performed over a range of x-positions
along the crack, yielding a crack opening profile along the crack, from
which the value at a 10 μ
m
distance from the crack tip was determined.
Due to the difficulty to accurately identify the crack tip location by
visual inspection systematically, a special criteria was employed to
assess the pixel coordinates of the crack tip in the images. Since the y-
coordinate of the crack is conveniently assessed by the above fit, i.e. as
the dparameter [43], the real difficulty was to assess the x-coordinate.
To this end, this location was defined as the x-position along the crack
at which the slope of the fitted step function, i.e. =
=bc|
df
dy yπ
02, falls
below a critical value taken as 0.1. This criteria is well motivated since
a too low value of the slope indicates that there is no discontinuity in
the displacement field. The particular value of 0.1 was chosen based on
agreement with manual inspections of crack tips in images.
In addition, based on the measured CTOD, the crack closure stress
was measured as the stress at which the CTOD exceeds 1 μ
m
. To in-
crease the reliability of this measurement, the average stress CTOD
curve over three subsequent cycles was considered. The reason for
taking this particular value was simply that it roughly corresponded to
the minimum detectable opening in view of the scatter over these three
cycles. The crack opening stress was only assessed at crack length of
three millimetres in all tests, which were used as a representative value
for the whole tests.
2.6. Modelling of the deformation behaviour at the crack tip
As a complement to the DIC measurement, finite element (FE)
modelling of the displacement field around the crack tip in the single
edge notched specimen was performed. The intention with this work
was to complement the DIC measurements as well as acquire more
detailed information about the local strain ahead of the crack tip and
the crack opening.
To this end, the same specimen model as used for the compliance
method and stress intensity factor computation was employed, except
for higher degree of mesh refinement at the crack tip corresponding to a
smallest element size of 5 μ
m
at the crack tip. However, rather than
restricting the behaviour to purely elastic, the material was given an
ideal plastic von Mises behaviour [39]. Hence, the effective von Mises
stress at yield, as well as the elastic modulus, was assessed as the
temperature dependent mechanical properties acquired from a series of
tensile tests at different temperatures performed by Rolls Royce plc.
This is a simplified approach but is motivated in view of the low rate of
hardening seen in tensile tests and the ensuing results in agreement
with experimental data as demonstrated later in Section 3.2. Moreover,
the employed temperature dependent thermal expansion coefficients
were calculated from the thermal strain measured in the TMF pre test
procedure.
The boundary conditions were applied in the same way as when
performing the compliance method described in Section 2.2. Similarly,
the same set of different models were set up, each having a plane crack
of different extension ranging from 0.5 to 5 mm measured from the
notch root. However, in order to represent the cyclic history associated
with the cyclic loading in the experiments, each model was cycled
between zero and 210 MPa over three cycles. The load values were
chosen to be the same as in the experiments, i.e. specimen
S
3
and
S5
in
Table 1. Additional cycles did not have any further significant effect
due to the ideal plastic material behaviour. Moreover, time dependent
uniform temperature fields were applied with the same cyclic variations
and magnitudes as in the experimental OP and IP cycling, i.e. between
400 °C and 750 °C.
In order to have consistency with the DIC analyses, the strain fields
acquired from the FE simulations were taken with reference to the
deformation state at zero nominal load of the same cycle, using avail-
able tools in Abaqus [39]. Furthermore, the computed thermal strain
was subtracted from this strain field output, in agreement with Eq. (1),
in order to acquire the mechanical strain field. A definition of CTOD
consistent with the DIC post processing was also employed, namely as
the interpolated node displacement in the y-direction at the position of
10 μ
m
from the crack tip location at the instant of maximum nominal
stress.
3. Results and discussion
3.1. Investigation of the effect of crack closure
In a wide context, crack closure has been demonstrated to affect the
fatigue crack growth rate in metallic materials [40,44]. Regrading the
present material type and load condition, the origin to crack closure has
been attributed to a number of phenomena including plasticity
[8,20,45,46], roughness [21] and oxide induced [7,12] crack closure.
However, without particular knowledge of the exact closure me-
chanism, recent studies [8,28–30] have demonstrated that variations in
crack growth rate caused by altering the load ratio and temperature
cycle are accounted for by compensating the stress intensity factor
range or the cyclic J-integral with respect to crack closure, i.e. Eq. (4) in
Section 2.2. For this reason, the effect of crack closure is investigated in
the present study to see whether it may explain the distinction between
in-phase (IP) and out-of-phase (OP) cycling.
Experimental determination of crack closure is however a subject of
controversy. In this investigation, two methods were employed, namely
by direct visual observation supported by digital image correlation
(DIC) and a method based on the measurement of the compliance
variation caused by crack closure, both explained in Sections 2.5 and
2.2 respectively. For the latter method, the compliance variation over a
given load cycle is converted to an opening parameter D, which by its
definition attains zero when completely closed and unity when com-
pletely open. Accordingly, the crack opening stress
σ
o
p
can then be
defined as a critical value of D, here selected as 0.9, see Section 2.3.
Fig. 3 compares the outcome when using the two different methods
to assess the degree of crack closure. In Fig. 3a, the crack tip opening
displacement (CTOD) measured using DIC is plotted as a function of
stress for one OP cycle and an IP cycle which occurred at a crack length
of about 3 mm in both cases. For the same OP and IP cycle, the crack
opening parameter Dis plotted as a function of stress in Fig. 3b. In the
former figure, crack opening is interpreted as the instant when CTOD
exceeds a limit value, chosen as 1 μ
m
, and in the former, when the
opening parameter exceeds 0.9, as motivated in Sections 2.5 and 2.2
respectively.
Interestingly, it is indicated that the two methods do not yield
consistent values of the crack opening stress
σ
o
p
. While the compliance
method results in a significant difference in crack opening stress be-
tween OP and IP, the DIC method asserts that the values are adjacent
and underestimated. The reason for this is mainly believed to be that
the critical value of the opening parameter 0.9, is not strict enough.
However, due to limitations of accuracy of measurement of the spe-
cimen compliance, it is meaningless to select a higher value, e.g. 0.95,
since such small deviation from unity is comparable in magnitude to the
measurement error of the opening parameter. Nevertheless, the DIC
method is argued to be more reliable since the method is based on the
local deformation events at the crack tip, rather than a macroscopic
variable such as the specimen compliance.
Using the crack opening stress acquired using the DIC method,
thermo mechanical faitgue (TMF) crack growth rate is correlated with
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
6
respect to the stress intensity factor and the effective stress intensity
factor for which crack closure is compensated as described in Section
2.2.6, see Fig. 4. Surprisingly, the variations seen in the IP tests of the
single edge notched (SEN) specimen in Fig. 4a are eliminated by
compensating for crack closure. A similar observation was made in the
investigation by Palmert et al. [8] in which isothermal and IP crack
growth tests with various load ratios and durations of dwell on a single
crystal nickel base alloy, collapsed into a single curve when adjusted for
crack closure. In contrast, the variation of growth rates seen in here is
likely to originate from the difference in the pre cracking procedure, in
view of the low amount of variation seen for the corner crack (CC)
specimen for which the pre crack procedure was not varied. Thus, it is
concluded that the pre cracking procedure, i.e. the choice of load ratio
and whether the pre crack cycling is thermo mechanical or conducted at
room temperature, does not explicitly affect the crack propagation rate.
Rather, it has an implicit effect which is entirely accounted for by the
influence of crack closure on fatigue crack propagation.
It is noted that for test conducted with the same pre-crack proce-
dure, the crack opening stress is lower in OP than in IP, see Fig. 3a. This
is in line with a recent study of a temperature dependent yield strip
model intended for TMF crack growth [45], which has indicated that
crack closure is more pronounced in IP compared to OP. However, it is
not evident why the difference in pre-cracking procedure has such a
significant effect on the crack opening level. It is suggested that the
different pre cracking procedures, see Table 1, have caused different
amount of residual plastic deformation in the notch which plausibly
could affect the degree of subsequent crack closure. Unfortunately, no
detailed information about the residual strain field is available, hence
this proposition cannot be confirmed at present.
Looking at the crack growth behaviour of both specimen types in
Fig. 4, it is seen that IP loading results in a distinctly higher crack
growth rate compared to OP loading. This is expected since in IP the
crack is exposed to the maximum load at the same instant of maximum
temperature, during which the material is less prone to withstand high
stresses and oxidation effects. Furthermore, based the results of the SEN
specimen, it is concluded that crack closure effect is of importance in
assessing the fatigue crack growth rate. However, it does not single
handedly account for the distinction between growth rates in OP and IP
loading.
3.2. Analysis of the crack tip deformation behaviour
A general aspect which may account for the differences between in-
phase (IP) and out-of-phase (OP) load conditions, is the stress-strain
state in the region ahead of the crack tip. Some authors have for in-
stance argued that inelastic creep deformation at the crack tip may
potentially contribute to increased crack growth rates in nickel base
superalloys [7,11,12,14,15,19,20,47]. Furthermore, it is also expected
that the crack tip stress strain state affects the diffusion of oxygen
[9,10,16,31],whose presence has been demonstrated to severely affect
the fatigue crack growth rate based on crack growth test under air and
vacuum conditions [9–17]. Thus, even in the absence of sustained loads
and an oxidising atmosphere, different thermo mechanical cycles may
impose different crack tip stress strain states due to the temperature-
dependence of the mechanical constitutive behaviour of the material.
This is an important aspect which must be clarified in order to fully
distinguish IP and OP load conditions.
Using digital image correlation (DIC) technique, the above pre-
sumption is affirmed, see Fig. 5a and b. The figure shows the me-
chanical strain field ahead of the crack tip at maximum applied stress of
210 MPa at one OP and one IP cycle which occurred at a crack length of
about 3 mm. Even though the differences are small, the IP cycle is seen
to impose a higher mechanical straining in the crack tip region. This is
further supported by the conducted finite element (FE) simulations
Fig. 3. (a) Crack tip opening displacement (CTOD) and (b) opening parameter Dderived from the specimen compliance, as a function of applied stress for one IP and
one OP cycle at a crack length of about 3 mm in both cases. The test parameters were 400–750 °C,
=
R
0
and =σ21
0
max MPa, and the tested specimens were
S3
and
S5
, presented in Table 1, for the OP and IP test respectively. In (a), the average, maximum and minimum value of CTOD over three subsequent load cycle.s are
displayed.
Fig. 4. Fatigue crack growth rate as a function of (a) stress-intensity factor and (b) effective stress intensity factor in the single edge notched (SEN) and corner cracked
(CC) specimen subjected to OP and IP TMF loading with 400–750 °C and
=
R
0
. All conducted TMF test are included, however they are only differentiated if the test is
OP or IP. The SEN specimen tested in IP have been subjected to different pre-crack procedures, as presented in Table 1. The effective stress intensity factor is
calculated according to Eq. (4) for which the stress opening stress is acquired using the DIC method, see Section 2.5.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
7
presented in Fig. 5c and d, which manifest a similar appearance in the
mechanical strain field, as well as a small difference between the OP
and IP cycle. Thus, there is consistent theoretical and experimental
support to argue that IP cycling results in higher deformation at the
crack tip region compared to OP cycling.
It should be remembered that the FE model in Figs. 5c and d does
not incorporate the effect of crack closure. Thus, the FE simulations
only reflect the difference originating from the effect of a phase shift in
the temperature cycle, without the eventual influence of crack closure.
Consequently, strain localisation at the crack tip occurs at all stress
levels in the FE model, even though it is not intuitively expected in
reality for stress levels below the crack opening stress. On the other
hand, crack closure occurs almost concurrently in the considered OP
and IP cycle, see Fig. 3a. Therefore, the difference seen in Fig. 5a and b
is mainly associated with the effect of the temperature cycle on the
thermo mechanical constitutive behaviour.
The difference in mechanical straining between IP and OP in ab-
sence of crack closure, is attributed to the temperature dependence of
the yield strength. Clearly, the material is at maximum temperature at
maximum stress, during which the yield strength is lower compared to
at minimum temperature. As a result, the crack tip is subjected to a
higher degree of inelastic deformation in contrast to an OP cycle, for
which inelastic deformation is better resisted thanks to the higher yield
strength at lower temperatures.
A larger crack tip deformation in the IP case is also reflected in the
measurement of the crack tip opening displacement (CTOD), see
Fig. 3a, in which CTOD is shown to be larger in IP compared to OP for
the same applied stress and crack length. This is further illustrated in
Fig. 6a, in which CTOD measured by DIC is plotted as function of crack
length. Effectively, the crack tip opens up more at maximum stress in IP
compared to OP for the major part of the tests.
In Fig. 6a, CTOD acquired from the FE simulations is included as
well, which demonstrates the overestimation of the experimentally
measured values. The discrepancy is again argued to be due to the
absence of crack closure in the FE model. To compensate for this, an
alternative definition of CTOD is attempted in which CTOD is assessed
as
=−
==
δδ δ||
FE σσ FE σσ
max o
p
(7)
where
δ
FE
is the closure-free CTOD computed in the FE model,
σ
max
is
the maximum stress of the cycle and
σ
o
p
is the crack opening stress
determined using DIC. This expression is a simple empirical estimate
similar to the proposition given by Donahue et al. [48], and accounts
for the effect of crack closure on CTOD by considering the crack
opening relative to the value at the crack opening stress
σ
o
p
. Based on
this measure, improved agreement with experimentally measured va-
lues is obtained, see Fig. 6b, even though CTOD in OP is still over-
estimated, which much likely is due to the absence of hardening be-
haviour in the FE model. Nevertheless, both experiments and the FE
simulations indicate that IP loading results in higher crack tip opening
compared to OP, also when the effect of crack closure is included.
Based on the above observation, it is well motivated to suggest that
the higher crack growth rate in IP, as demonstrated in Fig. 4b, may
originate from more severe mechanical load conditions at the crack tip
region. To validate this proposition, it is argued that CTOD is a suitable
parameter to relate to crack tip deformation, since theoretically, a
Fig. 5. Mechanical strain field in the tensile
direction measured using DIC in (a) an OP
and (b) IP cycle, as well as simulated by finite
element (FE) analysis, (c) and (d) for the OP
and IP cycle respectively. The mechanical
strain is measured at the maximum cyclic
stress with reference to the minimum cycle
stress of zero MPa at the cycle corresponding
to a crack length of 3 mm in both the ex-
periments and FE simulations. The test para-
meters were 400–750 °C,
=
R
0
and
=σ21
0
max MPa, and the studied specimens
were
S3
and
S5
, presented in Table 1, for the
OP and IP test respectively. In (a) and (b), the
mechanical strain field at maximum stress is
averaged over three subsequent cycles during
which the movement of the crack tip location,
as marked out by the circle, is negligible. To
ease the distinction between the figures, level
curves corresponding to the uniaxial yield
strain, which is about 0.5% over the tem-
perature range 400–750 °C, and 0.75%, i.e.
=≅εσE/0.5
yy y
% in black and =ε0.75
yy %in
red, are included. (For interpretation of the
references to color in this figure legend, the
reader is referred to the web version of this article.)
Fig. 6. Crack tip opening displacement
(CTOD) as a function of crack length
measured by DIC and modelled using FE
and Eq. (7). In (a) the experimentally
measured CTOD is compared with CTOD
measured in the FE directly, see Section
2.6, while in (b), the measured CTOD is
compared to the estimate given by Eq. (7).
The test parameters were 400–750 °C,
=
R
0
and =σ21
0
max MPa, and the tested
specimens were
S3
and
S5
, presented in
Table 1, for the OP and IP test respectively.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
8
higher CTOD can only be accommodated through higher straining of
the region ahead of the crack tip. The argument is also supported by the
investigation by Eckmann and Schweizer [20], in which damage evo-
lution ahead of the crack is correlated with the crack mouth opening,
which is similar parameter related to how much a crack is open.
Accepting this premise, the correlation between crack growth rate
and CTOD is made as similar to previous investigations [38,48–50], see
Fig. 7. Intriguingly, this correlation demonstrates the approaching of
the IP and OP curves compared to Fig. 4b, which suggests that the fa-
tigue crack growth rate in OP and IP are similar when compared at
equivalent mechanical conditions at the crack tip. In other words, the
crack growth rate seems to be dominated by the amount of crack tip
deformation and opening which is caused during a given cycle, in-
cluding the effects of crack closure and the constitutive behaviour of the
material.
The above demonstration suggests that oxidation plays a minor role
in distinguishing the mechanisms of crack growth in OP and IP loading,
in contrast to previous suggestions regarding isothermal dwell fatigue
[18] and TMF conditions [6,7]. However, as discussed in the next
section, oxygen is known to accelerate the crack growth rate at elevated
temperatures. Therefore, it is emphasised that crack growth is not in-
dependent of an oxidising atmosphere. Rather, it is proposed that the
difference in the effect of oxygen between IP and OP loading is small for
the tested material, as further discussed in the next section.
3.3. Metallographic investigation of the crack tip region
The results of the previous sections suggest that the effective in-
phase (IP) and out of phase (OP) TMF crack growth in RR1000 is mainly
controlled by the deformation caused at the crack tip. On the other
hand, many investigators have pointed out that crack tip oxidation
clearly influences the crack growth rate at constant elevated tempera-
ture for nickel base superalloys. In particular, it has been convincingly
demonstrated that the effect of oxygen accelerates the isothermal crack
growth rate by performing tests at different partial pressures of oxygen
[9–17], including investigations on RR1000 [11,14,16]. Furthermore,
grain boundary oxidation ahead of the crack tip has been reported in
Inconel 718 [19], as well as in RR1000 [14,22,23], under sustained
load conditions. Clearly, there is a potential influence of material re-
lated aspects, such as the influence of the microstructure and oxidation
at the high stressed crack tip region, which requires attention in view of
the results of the previous sections.
Fig. 8 displays the cross sectional appearance of the primary crack
tip in the centre of the specimen, i.e. middle position of the crack front,
for the OP and IP test interrupted at a crack length of roughly 4.2 mm in
both specimens. The figure also includes energy dispersive X ray
spectroscopy (EDS) results of the same area with respect the chemical
presence of oxygen. Clearly, oxygen is present in the crack as expected
in both the OP and IP specimen. However, the analysis, including the
visual inspection of the SEM images, does not indicate a significant
distinction regarding the role of oxides between the two crack tips.
These observations are in agreement with previous investigations of
fatigue crack tip oxidation in RR1000 subjected to sustained loads of
K
Imax,equal to 20
M
Pa m
during 10 min [51,52], and of K
Δ
equal to
17 and 30
M
Pa m
during one hour [18] at 700 °C in air, for which no
oxygen were identified ahead of the crack tip. On the other hand, longer
Fig. 7. Fatigue crack growth rate in the single edge notched specimen as a
function of CTOD measured using DIC, see Section 2.5, when subjected to TMF
cycling for which the test parameters were 400–750 °C,
=
R
0
and
=σ21
0
max MPa.
Fig. 8. Backscattered scanning elec-
tron microscopy (SEM) images of the
primary crack tip at a crack length of
roughly 4.2 mm in the specimen sub-
jected to (a) OP and (b) IP, as well as
EDS analyses of the same areas with
respect to chemical presence of oxygen
in (c) OP and (d) IP. The SEM images
are captured using 10 kV acceleration
voltage, 9 mm working distance and
with 10 kmagnification.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
9
dwell times have shown to yield traces of isolated oxides ahead of the
crack tip in RR1000 [14,22,23]. In the present study, the specimens
were only subjected to temperature in the range of 700–750 °C during
10 s, coincident with maximum load only for the IP tests. Thus, in view
of the tested thermo mechanical load cycle of 70 s cycle period, it is not
unexpected that the time is insufficient to cause significant crack tip
oxidation. This point clearly favours the argument that the assistance of
oxidation does not account for the difference between IP and OP crack
growth rates in RR1000, as suggested in the previous section.
On the other hand, the above argument is opposed by the observed
difference in crack path morphology between OP and IP crack growth.
In agreement to what is often reported for nickel base superalloys
[6,7,24], the conducted IP tests tend to manifest intergranular growth
while OP is more transgranular, see Fig. 9, which is consistent with the
general observation of increased tendency for intergranular growth
with increasing temperatures [15,25–27] and dwell times [14,15,53]
under isothermal conditions. More importantly, it is also reported that
the propensity of intergranular growth morphology decreases with
decreasing partial pressure of oxygen [9,11,12,14], which indicates a
dependency of the crack growth mechanism on the presence of oxygen.
Thus, based on the crack path morphology, the influence of oxygen on
the distinction between IP an OP crack growth in RR1000 cannot be
completely ruled out at this point.
Another observable feature in Fig. 9 is the difference in the caused
deformation in the wake of the growing crack. Based on the achieved
electron channelling contrast [18,54], the IP case demonstrates low
amount of deformation, except for the few grains which have failed
trangranularly, see Fig. 9b, whereas the OP case indicates severe de-
formation, see Fig. 9a. Thus, it appears as if plastic deformation asso-
ciated with the crack tip is accommodated by plastic deformation
within the grain in OP crack growth, while deformation in terms of
grain boundary sliding seems more justified in IP crack growth. By
some investigators [55,56], grain boundary sliding ahead of the crack
tip has been suggested to complement environmentally assisted inter-
granular cracking of nickel base superalloys. However, it has also re-
cently been demonstrated that grain boundary sliding occurs even
under vacuum, above 700 °C when subjected to tensile and creep load
conditions [57–60]. Thus, this indicates that oxygen is not necessarily
needed to cause intergranular deformation at elevated temperatures.
Consequently, it is possible that the difference in crack path mor-
phology between OP and IP is a result of changes in the deformation
behaviour of the microstructure between 400 °C and 750 °C, regardless
of the role of oxygen.
4. Conclusions
•Based on the adjacent correlations of both IP and OP crack growth
rates to the crack tip opening displacement, and the observed ab-
sence of oxygen ahead of the crack tip, it is proposed that the
difference between IP and OP TMF crack growth rate in coarse
grained RR1000 originates from the mechanical conditions at the
crack tip region, rather than environmental effects. This is supported
by digital image correlation analyses of the mechanical strain field
at the crack tip, which demonstrate that IP cycling leads to larger
crack tip deformation compared to OP at the same crack length and
load parameters (
=
σ
210
max MPa, =
R
0and
T
= 400–750 °C).
Moreover, finite element modelling suggests that the observed dif-
ferences in crack tip deformation and opening arise from the effect
of crack closure and the temperature dependence of the constitutive
behaviour of the material.
•For the tested single edge notched specimen, crack closure assessed
in terms of a crack opening stress
σ
o
p
, is highly dependent of the pre
cracking procedure, i.e. the choice of load ratio and whether the pre
crack cycling is thermo mechanical or conducted at room tem-
perature. However, the effective crack growth rate adjusted for
crack closure, is independent of pre cracking procedure.
Declaration of Competing Interest
The authors declare that they have no known competing financial
interests or personal relationships that could have appeared to influ-
ence the work reported in this paper.
Acknowledgement
This project has received funding from the European Union’s
Horizon 2020 research and innovation programme and Joint
Undertaking Clean Sky 2 under grant agreement No. 686600.
References
[1] Advisory council for aviation research and innovation in Europe.
[2] KoffBL, Fellow A, Gardens PB. Gas turbine technology evolution: a designers per-
spective. J Propul Power 2004;20(4):577–95.
[3] Metallic materials —Fatigue testing —Strain-controlled thermomechanical fatigue
testing method, Tech. rep.; 2011.
[4] McKnight RL, Laflen JH, Spamer GT. Turbine blade tip durability analysis. Tech.
rep., CR 165268. National Aeronautics and Space Administration; 1982.
[5] Sehitoglu H. Thermal and thermomechanical fatigue of structural alloys. Fatigue
Fract 1996;19:527–56.
[6] Pretty C, Whittaker M, Williams S. Crack growth of a polycrystalline nickel alloy
under TMF loading. Adv Mater Res 2014;891–892:1302–7.
[7] Pretty C, Whitaker M, Williams S. Thermo-mechanical fatigue crack growth of
RR1000. Materials 2017;10(1):34.
[8] Palmert F, Moverare J, Gustafsson D. Thermomechanical fatigue crack growth in a
single crystal nickel base superalloy. Int J Fatigue 2019;122:184–98.
[9] Andrieu E, Molins R, Ghonem H, Pineau A. Intergranular crack tip oxidation me-
chanism in a nickel-based superalloy. Mater Sci Eng A 1992;154(1):21–8.
[10] Pfaendtner J, McMahon C. Oxygen-induced intergranular cracking of a Ni-based
alloy at elevated temeratures - an example of dynamic embrittlement. Acta Mater
2001;49:3369–77.
[11] Knowles DM, Hunt DW. The influence of microstructure and environment on the
crack growth behavior of powder metallurgy nickel superalloy RR1000. Metall
Fig. 9. Backscattered scanning electron microscopy (SEM) images of the crack path in (a) OP, showing transgranular growth, and (b) IP, showing mixed inter and
transgranular morphology. The SEM images are captured using 8 kV acceleration voltage, 9 mm working distance and with 500 magnification.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
10
Mater Trans A: Phys Metall Mater Sci 2002;33(10):3165–72.
[12] Osinkolu GA, Onofrio G, Marchionni M. Fatigue crack growth in polycrystalline IN
718 superalloy. Mater Sci Eng A 2003;356(1–2):425–33.
[13] Winstone MR, Brooks JW. Advanced high temperature materials: aeroengine fa-
tigue. High Temp Fatigue 2008;20(1–2):15–24.
[14] Li HY, Sun JF, Hardy MC, Evans HE, Williams SJ, Doel TJ, et al. Effects of micro-
structure on high temperature dwell fatigue crack growth in a coarse grain PM
nickel based superalloy. Acta Mater 2015;90:355–69.
[15] Jiang R, Everitt S, Gao N, Soady K, Brooks JW, Reed PA. Influence of oxidation on
fatigue crack initiation and propagation in turbine disc alloy N18. Int J Fatigue
2015;75:89–99.
[16] Jiang R, Reed PA. Critical assessment 21: oxygen-assisted fatigue crack propagation
in turbine disc superalloys. Mater Sci Technol (United Kingdom) 2016;32(5):401–6.
[17] Jiang R, Bull DJ, Proprentner D, Shollock B, Reed PA. Effects of oxygen-related
damage on dwell-fatigue crack propagation in a P/M Ni-based superalloy: from 2D
to 3D assessment. Int J Fatigue 2017;99:175–86.
[18] Yu S, Li H, Hardy M, McDonald S, Bowen P. Mechanisms of dwell fatigue crack
growth in an advanced nickel disc alloy RR1000. MATEC Web of Conf
2014;14:03002.
[19] Gustafsson D, Moverare JJ, Johansson S, Simonsson K, Hörnqvist M, Mnsson T,
et al. Influence of high temperature hold times on the fatigue crack propagation in
Inconel 718. Int J Fatigue 2011;33(11):1461–9.
[20] Eckmann S, Schweizer C. Characterization of fatigue crack growth, damage me-
chanisms and damage evolution of the nickel-based superalloys MAR-M247 CC
(HIP) and CM-247 LC under thermomechanical fatigue loading using in situ optical
microscopy. Int J Fatigue 2017;99:235–41.
[21] McCarver JF, Ritchie RO. Fatigue crack propagation thresholds for long and short
cracks in René 95 Nickel-base superalloy. Mater Sci Eng 1982;55(1):63–7.
[22] Kitaguchi HS, Li HY, Evans HE, Ding RG, Jones IP, Baxter G, et al. Oxidation ahead
of a crack tip in an advanced Ni-based superalloy. Acta Mater 2013;61(6):1968–81.
[23] Kitaguchi HS, Moody MP, Li HY, Evans HE, Hardy MC, Lozano-Perez S. An atom
probe tomography study of the oxide-metal interface of an oxide intrusion ahead of
a crack in a polycrystalline Ni-based superalloy. Scripta Mater 2015;97:41–4.
[24] Hyde CJ, Sun W, Hyde TH. An investigation of the failure mechanisms in high
temperature materials subjected to isothermal and anisothermal fatigue and creep
conditions. Procedia Eng 2011;10:1157–62.
[25] Daus F, Li HY, Baxter G, Bray S, Bowen P. Mechanical and microstructural assess-
ments of RR1000 to IN718 inertia welds –effects of welding parameters. Mater Sci
Technol 2007;23(12):1424–32.
[26] Jiang R, Everitt S, Lewandowski M, Gao N, Reed PA. Grain size effects in a Ni-based
turbine disc alloy in the time and cycle dependent crack growth regimes. Int J
Fatigue 2014;62:217–27.
[27] Xu C, Yao ZH, Dong JX. The sharp drop in fatigue crack growth life at a critical
elevated temperature for a PM Ni-based superalloy FGH97. Mater Sci Eng A
2019;761(February).
[28] Jacobsson L, Persson C, Melin S. Thermo-mechanical fatigue crack propagation
experiments in Inconel 718. Int J Fatigue 2009;31(8–9):1318–26.
[29] Ewest D, Almroth P, Leidermark D, Simonsson K, Sjödin B. Fatigue crack propa-
gation in a ductile superalloy at room temperature and extensive cyclic plastic flow.
Int J Fatigue 2015;80:40–9.
[30] Ewest D, Almroth P, Sjödin B, Leidermark D, Simonsson K. Isothermal and ther-
momechanical fatigue crack propagation in both virgin and thermally aged Haynes
230. Int J Fatigue 2019;120(October 2018):96–106.
[31] Evans HE, Li HY, Bowen P. A mechanism for stress-aided grain boundary oxidation
ahead of cracks. Scripta Mater 2013;69(2):179–82.
[32] McEvily AJ, Gonzalez Velazquez J. Fatigue crack tip deformation. Metall Trans A
1992;23(8):2211–21.
[33] Hardy M, Zirbel B, Shen G, Shenkar R. Developing damage tolerance and creep
resistance in a high strength nickel alloy for disc applications. Superalloys
2004;2004:83–90.
[34] Mitchell R, Lemsky J, Ramanathan R, Li H, Perkins K, Connor L. Process develop-
ment and microstructure and mechanical property evaluation of a dual micro-
structure heat treated advanced nickel disc alloy. Superalloys 2008;2008:347–56.
[35] Mitchell R, Hardy M, Preuss M, Tin S. Development of γmorphology in P/M rotor
disc alloys during heat treatment. Superalloys 2004;2004:361–70.
[36] Collins DM, Heenan RK, Stone HJ. Characterization of gamma prime (γ) precipitates
in a polycrystalline nickel-base superalloy using small-angle neutron scattering.
Metall Mater Trans A 2011;42(1):49–59.
[37] Ewest D, Almroth P, Sjödin B, Simonsson K, Leidermark D, Moverare J. A modified
compliance method for fatigue crack propagation applied on a single edge notch
specimen. Int J Fatigue 2016;92:61–70.
[38] Norman V, Skoglund P, Leidermark D, Moverare J. The transition from micro- to
macrocrack growth in compacted graphite iron subjected to thermo-mechanical
fatigue. Eng Fract Mech 2017;186:268–82.
[39] Abaqus, Abaqus/CAE user’s manual, ver. 6.12; 2012.
[40] Elber W. The significance of fatigue crack closure. Damage tolerance in aircraft
structures. ASTM International; 1971.
[41] ASTM, Standard Test Method for Measurement of Fatigue Crack Growth Rates,
Tech. rep.; 2014.
[42] Eberl C, Thompson R, Gianola D. Matlab-based DIC code, version 2. Matlab’s File
Exchange; 2015.
[43] Vasco-Olmo JM, Díaz FA, Antunes FV, James MN. Experimental evaluation of CTOD
in constant amplitude fatigue crack growth from crack tip displacement fields.
Frattura ed Integrita Strutturale 2017;11(41):157–65.
[44] Suresh S. Fatigue of materials, second ed. Cambridge: Cambridge University Press;
1998.
[45] Fischer C, Schweizer C, Seifert T. Assessment of fatigue crack closure under in-phase
and out-of-phase thermomechanical fatigue loading using a temperature dependent
strip yield model. Int J Fatigue 2015;78:22–30.
[46] Fischer C, Schweizer C, Seifert T. A crack opening stress equation for in-phase and
out-of-phase thermomechanical fatigue loading. Int J Fatigue 2016;88:178–84.
[47] Addison DA, Tucker JD, Siegmund T, Tomar V, Kruzic JJ. Cyclic and time-depen-
dent crack growth mechanisms in Alloy 617 at 800 C. Mater Sci Eng A
2018;737(July):205–12.
[48] Donahue RJ, Clark HM, Atanmo P, Kumble R, McEvily AJ. Crack opening dis-
placement and the rate of fatigue crack growth. Int J Fract Mech 1972;8(2):209–19.
[49] Liu HW. Fatigue crack growth by crack tip cyclie plastic deformation: the unzipping
model. Int J Fract 1989;39(1–3):63–77.
[50] Antunes FV, Branco R, Prates PA, Borrego L. Fatigue crack growth modelling based
on CTOD for the 7050–T6 alloy. Fatigue Fract Eng Mater Struct
2017;40(8):1309–20.
[51] Viskari L, Hörnqvist M, Moore KL, Cao Y, Stiller K. Intergranular crack tip oxidation
in a Ni-base superalloy. Acta Mater 2013;61(10):3630–9.
[52] Hörnqvist M, Viskari L, Moore KL, Stiller K. High-temperature crack growth in a Ni-
base superalloy during sustained load. Mater Sci Eng A 2014;609:131–40.
[53] Moverare JJ, Gustafsson D. Hold-time effect on the thermo-mechanical fatigue
crack growth behaviour of Inconel 718. Mater Sci Eng A
2011;528(29–30):8660–70.
[54] Gutierrez-Urrutia I, Zaefferer S, Raabe D. Coupling of electron channeling with
EBSD: toward the quantitative characterization of deformation structures in the
sem. Jom 2013;65(9):1229–36.
[55] Chen X, Yang Z, Sokolov MA, Erdman DL, Mo K, Stubbins JF. Effect of creep and
oxidation on reduced fatigue life of Ni-based alloy 617 at 850 °c. J Nucl Mater
2014;444(1–3):393–403.
[56] Németh AA, Crudden DJ, Armstrong DE, Collins DM, Li K, Wilkinson AJ, et al.
Environmentally-assisted grain boundary attack as a mechanism of embrittlement
in a nickel-based superalloy. Acta Mater 2017;126:361–71.
[57] Soula A, Renollet Y, Boivin D, Pouchou JL, Locq D, Caron P, et al. Analysis of high-
temperature creep deformation in a polycrystalline nickel-base superalloy. Mater
Sci Eng A 2009;510–511:301–6.
[58] Walley JL, Wheeler R, Uchic MD, Mills MJ. In-situ mechanical testing for char-
acterizing strain localization during deformation at elevated temperatures. Exp
Mech 2012;52(4):405–16.
[59] Carter JL, Kuper MW, Uchic MD, Mills MJ. Characterization of localized deforma-
tion near grain boundaries of superalloy René-104 at elevated temperature. Mater
Sci Eng A 2014;605:127–36.
[60] Pataky GJ, Sehitoglu H. Experimental methodology for studying strain hetero-
geneity with microstructural data from high temperature deformation. Exp Mech
2015;55(1):53–63.
V. Norman, et al. International Journal of Fatigue 135 (2020) 105528
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