Creep-fatigue of High Temperature Materials for VHTR: Effect of ...

INL/CON-11-20780
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Creep-Fatigue of High
Temperature Materials
for VHTR: Effect of
Cyclic Loading and
Environment
ICAPP 2011
C. Cabet
L. Carroll
R. Madland
R. Wright
May 2011
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Creep-fatigue of High Temperature Materials for VHTR:
Effect of Cyclic Loading and Environment
Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
C. Cabet
CEA, DEN, DPC, SCCME, Laboratoire d'Etude de la Corrosion Non Aqueuse, F-91191 Gif-sur-Yvette, France
Tel: +33 169 081 615, Fax: +33 169 081 586, Email: celine.cabet@cea.fr
L. Carroll
Idaho National Laboratory
Idaho Falls, ID, USA
R. Madland
Colorado School of Mines
Denver, CO, USA
R. Wright
Idaho National Laboratory
Idaho Falls, ID, USA
Abstract – Alloy 617 is the one of the leading candidate materials for Intermediate Heat
Exchangers (IHX) of a Very High Temperature Reactor (VHTR). System start-ups and shut-downs
as well as power transients will produce low cycle fatigue (LCF) loadings of components.
Furthermore, the anticipated IHX operating temperature, up to 950°C, is in the creep regime so
that creep-fatigue interaction, which can significantly increase the fatigue crack growth, may be
one of the primary IHX damage modes. To address the needs for Alloy 617 codification and
licensing, a significant creep-fatigue testing program is underway at Idaho National Laboratory.
Strain controlled LCF tests including hold times up to 1800s at maximum tensile strain were
conducted at total strain range of 0.3% and 0.6% in air at 950°C. Creep-fatigue testing was also
performed in a simulated VHTR impure helium coolant for selected experimental conditions. The
creep-fatigue tests resulted in failure times up to 1000 hrs. Fatigue resistance was significantly
decreased when a hold time was added at peak stress and when the total strain was increased. The
fracture mode also changed from transgranular to intergranular with introduction of a tensile
hold. Changes in the microstructure were methodically characterized. A combined effect of
temperature, cyclic and static loading and environment was evidenced in the targeted operating
conditions of the IHX. This paper reviews the data previously published by Carroll and coworkers
in references 10 and 11 focusing on the role of inelastic strain accumulation and of
oxidation in the initiation and propagation of surface fatigue cracks.
I. INTRODUCTION
Alloy 617 is the leading candidate material for
Intermediate Heat eXchangers (IHX) of helium-cooled
Very High Temperature Reactor (VHTR) systems.
Conceptual design requires an outlet temperature of greater
than 850°C to efficiently co-generate hydrogen and
electricity, with a maximum expected temperature of
950°C 1 . The IHX will operate at the reactor outlet
temperature of up to 950°C.
Reactor start-ups and shut-downs as well as power
transients will produce low cycle fatigue (LCF) loadings of
components. Furthermore, the anticipated IHX operating
temperature is in the range where creep deformation occurs
so that interaction between a growing fatigue crack and the
bulk creep damage, which can significantly increase the
312
fatigue crack growth rate, may be one of the primary IHX
damage modes. However, in the draft code case that was
developed in the 1980s to have Alloy 617 approved in
ASME Code section III for nuclear use at high
temperature 2 , creep-fatigue was not sufficiently addressed.
This calls for a more complete database for creep-fatigue
behavior of Alloy 617; for a better understanding of the
coupled influence of aging, fatigue, creep and corrosion;
and for a better life prediction of components under creepfatigue
deformation. �
A significant testing program is underway at Idaho
National Laboratory to address the needs for codification
and licensing. To reproduce the expected deformation in a
laboratory setting, creep-fatigue testing introduces a hold
time in a strain-controlled fatigue cycle. Previous work on
Alloy 617 has suggested that a hold time during the tensile

portion of the fatigue cycle is more damaging 3-5 than
compressive holds or both a tensile and compressive hold.
Therefore holds at peak tensile strain were investigated.
The helium coolant in the primary circuit of a VHTR
system is expected to contain low levels of impurities 6 such
as H2, H2O, CO, CO2, CH4, etc, that can react toward
metallic materials at high temperature. As far as corrosion
of Alloy 617 is concerned, the optimum coolant chemistry
will enable the alloy to be oxidized in service for the
formation of a surface, continuous and dense, chromia
scale which would protect the bulk alloy and has little
influence on the mechanical properties 7 . This corrosion
behavior is similar to oxidation in air at the same
temperature 8,9 . Hence, experiments were performed in air
as well as in controlled VHTR-simulated helium to account
for any possible environmental effect. To enhance the
understanding of creep-fatigue deformation mechanisms at
950°C, the fatigue and creep-fatigue behavior of Alloy 617
is presented with systematic microstructure characterization
of the alloy surface, the specimen bulk and the vicinity of
the cracks. This paper reviews the creep-fatigue data
presented in two previous references by Carroll et al.
10, 11
with an emphasis on the role of inelastic strain
accumulation and the effect of oxidation in the initiation
and propagation of surface fatigue cracks.
II. EXPERIMENTAL
II.A. Material
Cylindrical creep-fatigue specimens, 7.5 mm diameter
in the reduced section and a gage length of 12 mm, were
machined from Alloy 617 annealed plate; the long axis of
the specimen aligned with the rolling direction. The
composition of the investigated heat of Alloy 617 is given
in Table I.
TABLE I
Alloy 617 chemical composition in wt.%.
Ni C Cr Co Mo Fe Al Ti Si Mn Cu
bal 0.08 21.9 11.4 9.3 1.7 1.0 0.3 0.1 0.1 0.04
II.B. Apparatus
Cyclic testing was conducted on servo-hydraulic test
machines in axial strain-control mode in accordance with
the ASTM Standard E606-04 12 . Radio-frequency induction
heating was used to heat the specimens. Temperature
control was achieved using a combination of spot-welded
thermocouples on a shoulder of the specimen and a
thermocouple loop at the center of the gage section. The
temperature gradient was measured with a specimen
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Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
instrumented with spot-welded thermocouples along the
gage length and was found to deviate by less than 1%
within the gage section.
For the cyclic tests in controlled chemistry helium, a
servo-hydraulic test frame equipped with a gas-tight
chamber was used. A gas system was developed to feed
VHTR simulated coolant to the chamber. It comprises a
manifold containing bottles of helium with controlled
partial pressures of CO, CH4, and H2; mass flow controller
to regulate the helium flow; and pressure controllers. Small
quantities of H2O are added to control the partial pressure
of water vapor in the helium flow (P(H2O) as low
as 5μbar). A gas chromatograph and two solid-state
hygrometers record the gas chemistry at the inlet and outlet
of the chamber. Feedback is employed on the inlet
hygrometry.
II.C. Testing methods
Fully reversed strain controlled low cycle fatigue and
creep-fatigue testing was completed on Alloy 617 at a total
strain range of 0.3% and 0.6% in air and a simulated
VHTR helium at 950°C. A triangular waveform and a ramp
rate of 10 -3 /s were used for low cycle fatigue testing.
Creep-fatigue testing followed a trapezoidal waveform with
a tensile hold imposed at the maximum tensile strain.
Tensile holds of 180, 600, and 1800 sec were investigated.
The number of cycles to failure, Nf, is defined as the point
at which the ratio of the peak tensile stress to the peak
compressive stress initially deviated from its level slope
and the point at which the ratio was 80% of the value at
deviation 12 . Test completion was prior to actual specimen
separation.
For testing in helium, a gas mixture was selected to
roughly simulate the controlled VHTR coolant. It contained
50 μbar CO and 350 μbar H2 in addition to approximately
8 μbar H2O at 1.1 bar total helium pressure to be oxidizing
with respect to Alloy 617.
II.D. Specimen characterization
After creep-fatigue deformation, the gage sections of
the deformed specimens were cut along the stress axis
through the largest surface crack. The specimen was then
coated with a thin layer of gold followed by nickel plating
to ensure retention of the surface oxide. The coated
specimens were mounted in phenolic, polished to a mirror
finish, and then etched with a 2% bromine solution in
methanol. Metallurgical evaluation was conducted with
optical microscopy and Field Emission Gun-Scanning
Electron Microscopy (FEG-SEM) equipped with an EDX
spectrometer.

III. RESULTS
III.A. Low cycle fatigue behavior
Continuous low cycle fatigue (LCF) testing was
completed at 950°C to provide a baseline for the creepfatigue
behavior. Two total strain ranges were investigated,
0.3% and 0.6%, in air and in helium. The number of cycles
to failure and the total test time are included in Table II.
TABLE II
Fatigue and creep-fatigue tests completed at 950°C and a
strain rate of 10 -3 /s.
Strain range
(%)
Hold time
(s)
Environment
Nf (cycle)
Time
(h)
0 air 9641 16
0 air 7133 12
0 air 5867 10
0 air 9000 15
0 helium 8000 13
0 helium 7333 12
180 air 3989 210
180 air 2485 130
0.3
180
180
air
helium
4486
3004
230
155
180 helium 3373 174
600 air 4096 690
600 air 2623 440
600 air 4361 735
600 air 4430 745
600 helium 3130 530
1800 air 4805 2410
1800 air 4650 2330
0 air 1722 6
0 air 1390 5
0 air 1480 5
0.6
180
180
air
air
950
922
51
49
600 air 686 117
600 air 634 108
1800 air 661 333
The peak tensile and compressive stresses observed in
the LCF tests at 0.3% and 0.6% total strain were relatively
symmetrical, i.e. the peak tensile and compressive stresses
were of the same magnitude, as shown in Figures 1a and
2a, respectively. Furthermore, a steady state stress was
reached in approximately 10 cycles and remained constant
until macrocrack initiation or just prior to failure. Also,
note that the steady state peak stress in the 0.6% total strain
test was similar in magnitude to the 0.3% total strain test.
This is consistent with the shape of the hysteresis loops
shown in Figures 1b and 2b for the 0.3% and 0.6% total
strain conditions, respectively. The hysteresis loops were
relatively unchanging for cycles 9, 99, and 999 and thus the
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Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
inelastic strain also did not change significantly as a
function of cycle.
III.B. Creep-fatigue behavior
Creep-fatigue testing was also conducted at 950°C and
at 0.3% and 0.6% total strains with tensile hold times of up
to 1800s. A list of the creep-fatigue conditions and fatigue
life is shown in Table II; total test times were as long as one
month.
The peak tensile and compressive stresses are also
shown as a function of cycle for the creep-fatigue tests
(Figures 1a and 2a). The creep-fatigue peak stresses versus
cycle profiles were similar regardless of the duration of the
tensile hold. The creep-fatigue peak stresses versus cycle
profiles were also relatively symmetric, although the
magnitude of the stresses in compression were slightly
greater than in tension. The peak stresses did not achieve a
steady state value, instead the peak stresses slowly
decreased with cycle and had two transition points
following which a more rapid decrease was observed. The
hysteresis loops shown in Figures 1c and 2c illustrate this
cyclic softening. The loops shown for cycles 999 and 499
demonstrated a lower peak stress magnitude versus cycle
99, although the width of the loops at zero stress, the
inelastic strain, did not vary greatly.
III.C. Creep-fatigue life in air and in helium
Figure 3 and Table II give the cycles to failure for LCF
and creep-fatigue testing at 950°C in air and in helium.
The continuous cycle fatigue lives were longer than
corresponding creep-fatigue lifetimes. For example in the
case of the 0.3% total strain range, the addition of a 180s
hold time reduced the number of cycles to failure by a
factor of 2. However, the increasing duration of tensile hold
did not further decrease the creep-fatigue life at least at a
total strain range of 0.3%, i.e. the cycle life with a 180s
hold was comparable to the cycle life with a 1800s hold.
The influence of the hold duration on the fatigue life at a
0.6% total strain still needs to be ascertained.
For continuous fatigue testing as well as for creepfatigue
at 950°C, the cycle life of Alloy 617 in air and in
VHTR simulated helium was within the same scatter band,
as shown in Figure 3. Similar peak tensile stresses were
also observed in air and impure helium (see Figure 4).
Again in continuous cycle fatigue, the peak stresses
achieved a steady state value while cyclic softening was
observed in creep-fatigue.

a.
b.
c.
Fig. 1. LCF and creep-fatigue tests in air at 950°C and a
0.3% total strain: a) peak tensile and compressive stresses as a
function of cycle; b) hysteresis loops for a LCF test and c)
hysteresis loops for a 60s hold creep-fatigue test.
315
a.
b.
c.
Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
Fig. 2. LCF and creep-fatigue tests in air at 950°C and a
0.6% total strain: a) peak tensile and compressive stresses as a
function of cycle; b) hysteresis loops for a LCF test and c)
hysteresis loops for a 60s hold creep-fatigue test.

Cycle to failure
10000
1000
100
0.3% total strain in air
0.3% total strain in helium
0.6% total strain in air
0,01 1 100 10000
Hold time (s)
Fig. 3. Cycles to failure as a function of hold time for LCF
and creep-fatigue tests in air and in helium at a 0.3% and 0.6%
total strain.
a.
b.
Fig. 4. Peak tensile stresses as a function of cycle in air and
in helium at a 0.3% total strain for a) LCF and cb) reep-fatigue
with a 180s hold time.
Typical stress relaxation curves at mid-life are shown
in Figure 5 for creep-fatigue tests with a hold time up to
1800s with a 0.3% and 0.6% total strain. During the hold
period at constant strain, the stress relaxed at the same rate
for all hold durations: it initially decreased rapidly and was
almost completely relaxed after approximately 180s into
the hold.
316
a.
b.
Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
Fig. 5. Stress relaxation at mid-life during the tensile hold
for creep-fatigue specimens tested in air a) at 0.3% total strain
with a 180s, 600s, and 1800s hold time and b) at 0.6% total strain
with a 180s, and 600s hold time.
III.D. Failure mode
No noticeable difference was observed in specimen
cracking after cyclic testing in air and in helium. Hence in
the following, results from both environments are presented
together.
In continuous cycling tests, the primary crack (most
likely initiated and) propagated in a predominantly
transgranular manner, perpendicular to the stress axis.
Figure 6 shows longitudinal cross sections through the
specimen gage after LCF. The edges of the cracks were
slightly oxidized and there was little indication of an oxide
having formed on the surface of the specimens. A few wide
but short intergranular secondary cracks were observed
perpendicular to the stress axis in the specimens tested at a
0.6% total strain.
The addition of a hold time resulted in multiple larger
or primary cracks originating from the specimen surface as
well as an evolving microstructure, particularly at the
specimen surface. For the specimens tested in creepfatigue,
all of the grain boundaries close to the specimen

surface and perpendicular to the stress axis were oxidized
(the intergranular oxide was mainly Al-rich). A majority of
these surface internal oxides have initiated intergranular
cracks (after opening, the edges of these small secondary
cracks have oxidized and eventually the cracks were fully
filled by the Cr-rich oxide) as can be seen in Figure 7a. It is
difficult to determine on the initiation mode of the primary
crack (or cracks) because of the evolving microstructure at
the surface and surrounding the cracks. However, based on
the many shorter secondary cracks which started at
oxidized grain boundaries, it is likely that the primary
cracking also initiated intergranularly.
a.
Fig. 6. Cracking in continuously cycled specimens deformed
in air at a) a 0.3% total strain and b) a 0.6% total strain. The
stress axis is horizontal and in the plane of the page.
The crack propagation was intergranular. In Figure 7a,
the primary crack was open and an oxide, rich in Cr and
containing Ti, has formed along the flanks. Grain
boundaries intersecting with the main crack exhibit fine
carbides and Al-rich precipitates, most certainly alumina.
Figure 7b illustrates oxidation at the crack tip. Alumina was
detected ahead of the main crack tip as fine intergranular
veins.
III.E. Surface evolution
The surface microstructure in the specimens cycled in
air was qualitatively similar to those cycled in impure
helium. Therefore, the results reported in this section do
not differentiate between testing environment. In addition
the microstructure of deformed specimens was consistent
with statically exposed specimens 8,13 with the exception of
occurrence of secondary and primary cracking.
b.
317
a.
b.
intergranular
crack
Al-rich
precipitates
Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
hole
Cr-rich oxide
40 µm
100 µm
Fig. 7. Cracking in creep-fatigue specimens deformed in air
at a 0.3% total strain with a) a 600s hold and b) a 180s hold
(close up of the crack tip). The stress axis is horizontal and in the
plane of the page.
Figure 7a exemplifies the basic features of the surface
evolution:
- A Cr-rich oxide scale has formed which covers the
entire surface. Variations in the scale thickness were
related to the total time exposed at 950°C: after
continuous cycle fatigue, the oxide was relatively
thin (approximately 2µm), while after creep-fatigue
deformation a thicker oxide had formed (up to
10µm).
- Internal Al-rich oxides, precipitated beneath the
chromia scale and at grain boundaries close to the
surface. In continuously cycled specimens, the
intergranular oxides were thin and finger-like. In
creep-fatigue deformed specimens, the depth of
grain boundary oxides was greater, and as
previously explained, some had evolved into
secondary cracks filled with a Cr-rich oxide.

- An area free of grain boundary carbides has evolved
in the subsurface. Such a carbide-depletion was also
evidenced in the region surrounding the primary
creep-fatigue cracks. Carbide dissolution was
related to the diffusion of chromium to the surface
(of the specimen or of the crack) to form the
chromia scale. Chromium removal then may have
destabilized the grain boundary Cr-rich carbides in
the subsurface area.
III.F. Bulk evolution
For pure fatigue specimens (and oxidized coupons
without loading as well), the material bulk was clean from
cracking. On the contrary, cracks were observed in the bulk
of the creep-fatigue tested specimens. Examples of these
cracks are shown in the images in Figure 8.
a.
b.
50 µm
50 µm
Fig. 8. Grain boundary damage observed in the bulk material
of specimens deformed at a 0.3% total strain range with a) a 180 s
hold in impure helium and b) a 600 s hold in air. The stress axis is
horizontal and in the plane of the page.
The cracks were relatively thin and typically followed
grain boundaries perpendicular to the stress axis. Their
length varied, from several micrometers to several grain
diameters in length. In some cases, there were boundaries
with multiple small cracks. Although this type of damage
was randomly distributed in the bulk, most of the grain
boundaries were completely free of cracking and these
intact grain boundaries did not exhibit cavitation and
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Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
instead microscopically appeared free of grain boundary
damage (at magnifications of up to 100,000 times). The
bulk cracks were absent of any oxide and thus it can be
concluded that they were not surface connected and
exposed to the environment.
IV. DISCUSSION
IV.A. Effect of an air or helium environment
The cycle lives of continuous cycle and creep-fatigue
of Alloy 617 in air and in VHTR simulated helium were
within the same scatter band at 950°C, as shown in
Figure 3. This tendency is rather consistent with the small
number of previous studies on the influence in fatigue of air
compared to impure helium. Nagato and coworkers 14
observed that the fatigue life of Hastelloy X, a nickel base
alloy rich in Cr, was longer in impure helium than in air but
the effect of environment was less remarkable at
temperatures above 800°C and for creep-fatigue. At
temperatures of 750°, 850° and 950°C, Meurer and
coworkers 5 observed a slight increase in the fatigue life of
Alloy 617 tested in impure helium compared to air at a
total strain of 0.3%; this difference was not evidenced at
higher strain ranges. Strizak and coworkers 15 observed that
a helium environment was not detrimental to fatigue life of
Alloy 617 at temperatures of up to 704°C and Rao and
coworkers 3 concluded that at 950°C the influence of
environment, if any, is strongly reduced for longer hold
times. In addition to an equivalent fatigue life, the
tensile/compressive peak stress versus cycle, the hysteresis
loop, and the stress relaxation during the peak tensile hold
were similar under both atmospheres (see Figure 4). No
difference in the crack initiation behavior, propagation
mode, or oxidation of the cracks was evidenced between
the air and impure helium cycled specimens. Finally,
metallurgical characterization of the specimens cyclically
deformed in air and in helium failed to reveal any
significant differences in the oxidation products or surface
morphology. The corrosion product development and the
surface evolution are consistent with the general corrosion
behavior of Alloy 617 under a high temperature oxidizing
environment as is air or controlled VHTR helium coolant
(helium with significant water vapor and carbon monoxide
partial pressures at 950°C). In an atmosphere with a
sufficient oxidation potential, Alloy 617 forms a chromia
surface scale and aluminum oxidizes internally 9,13 . These
similarities in the overall behavior of the specimens cycled
in air and in simulated impure VHTR helium allow for the
joint discussion of the results for both environments
without specific reference to the testing atmosphere.

IV.B. Basis for creep-fatigue interaction
Creep-fatigue interaction is a combination of creep
damage and fatigue damage which accelerates the
deformation process. Typically three types of failure are
defined: fatigue-dominated, creep-dominated, and creepfatigue
interaction 16-18 . Fatigue dominated failures are
typically observed at lower to intermediate temperatures
and transgranular crack initiation and propagation is
observed. On the contrary, creep dominated failures occur
at high temperatures and intergranular crack initiation and
propagation with extensive creep cavitation occurs. Creepfatigue
interaction is illustrated as mixed-mode crack
propagation and the presence of creep cavitation on the
grain boundaries in the bulk material. Cavity formation and
fatigue crack initiation and propagation independently
develop, then interaction occurs if one mode accelerates the
other damage process. The creep and fatigue failure modes
interact through cavitation accelerating the crack initiation
or propagation process or fatigue deformation enhancing
cavitation 18 . Identifying the dominant failure mode and the
influence of environment is important for life prediction
and modeling efforts.
On the one hand, the creep-fatigue specimens did not
fail in a fatigue-dominated manner at 950°C. The primary
crack in the continuously cycled specimens initiated and
propagated in a transgranular manner (fatigue dominated
manner). Oxidized grain boundaries were occasionally
present but did not result in specimen failure. For all cases
in creep-fatigue, several longer cracks likely initiated and
did propagate intergranularly. Rao and coworkers 3 also
reported transgranular cracking of Alloy 617 at 950°C in
fatigue cycled specimens and the addition of a 60s tensile
hold resulted in the cracking initiating transgranularly and
propagating by a mixed mode. Besides, at lower strain rates
in a helium environment: approximately 10 -4 /s, the crack
propagation mode switched to intergranular.
On the other hand, massive cavitation which would
have caused creep dominated failure was not observed in
any of the creep-fatigue specimens deformed at 950°C with
holds as long as 1800s. Grain boundary cavities were
depicted for many alloys in the regime of creep-fatigue
interaction 3,16,18-20 . A detailed assessment correlated the
size and number of grain boundary cavities and the
accumulated creep damage for 316 stainless steel 16 . It was
also found that a critical amount of deformation was
required before creep cavities nucleated. Recent work by
Lillo and co-workers 21 suggested that significant creep
cavitation did not occur in Alloy 617 until greater than
10% creep strain during creep deformation in the
temperature range from 900°C to 1000°C. Inelastic strain
was evaluated by integrating the stress relaxation data from
Figure 5 [this calculation was done by S. Sham from Oak
Ridge National Laboratory]. The total inelastic strain was
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Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
calculated by integrating the stress relaxation curve at a
mid-life cycle. The accumulated creep strain was estimated
to be less than 2.5% and 0.5% for respectively the 0.3%
and the 0.6% total strain tests,. Therefore, creep cavities
are not expected to be formed in the present testing of
Alloy 617 at 950°C.
IV.C. Effect of a hold period in tensile strain
Similar to what has been observed for stainless steels
and nickel base alloys, the introduction of hold times at
peak tensile strain in continuous cycle fatigue reduces the
fatigue life of Alloy 617 (Figure 3). For example at a 0.3%
total strain, the introduction of a 180s hold period at peak
tensile strain decreased the life by approximately a factor
of 2. A greater decrease in cycles to failure at lower strain
ranges was also observed, consistent with the results on
Nimonic PE-16, a Ni-Cr superalloy 19 .
The drop in cycle life is likely associated with the
change in the cracking mode. The shift in crack
propagation mode from transgranular to intergranular may
be a result of grain boundary damage from creep processes
that accelerates intergranular crack propagation (see
paragraph IV.B) but may be also influenced by the
environment accelerating grain boundary crack propagation
and promoting intergranular failure.
Interestingly, the increasing duration of the peak
tensile hold at 0.3% total strain did not further decrease the
number of cycles to failure; that is to say, the 180s hold
creep-fatigue life was similar to that of the 1800s hold. The
amount of creep damage, estimated by the accumulated
inelastic strain, is relatively similar with increasing hold
periods as the stresses during the tensile hold were almost
completely relaxed by the end of the first 180s.
This suggests creep-fatigue interaction in the current
testing. Although creep intergranular cavities were not
observed, bulk cracking was evidenced in all of the creepfatigue
deformed specimens as illustrated in Figure 8.
Because intergranular crack was not found in the
continuously cycle fatigue specimens (or in unloaded
material), it was assumed that bulk cracking was caused by
creep deformation.
The amount of bulk material cracking was quantified
to determine if a relationship existed between the duration
of the hold time and the amount of grain boundary damage.
The total number of cracks, total length of cracks, and the
average length of the cracks were determined for a
(0.5mm)² area at five locations, as shown schematically in
Figure 9.
The areas analyzed were 1mm in each direction (in the
plane of the page) from the primary crack tip and 3mm in
each direction along the stress axis (in the plane of the
page) from the primary crack tip. For all specimens, the
cracking was statistically similar among the five areas and
no general trends were observed. In particular, the amount

or length of the bulk cracking did not appear to be related
to location relative to the primary crack. Hence, the
cracking features (number, average length, cumulated
length of the cracks) were calculated for the all of the areas
combined. Figure 10 shows the results graphically as a
function of hold time. As mentioned previously, there was
no cracking observed in the continuous cycle fatigue
specimens. It was also evidenced that increasing hold times
induced relatively similar amounts of grain boundary
cracking. Additionally, whatever the environment, air or
impure helium, the material exhibited a similar degree of
bulk grain boundary cracking.
Fig. 9. Schematic of the statistical method used to quantify
damage in the LCF and creep-fatigue deformed specimens. The
stress axis is horizontal and in the plane of the page.
It is believed that creep-fatigue interaction may be
caused by bulk cracking accelerating the surface fatigue
crack propagation. Equivalent creep strain accumulated for
a 180s hold and a 1800s hold (because of the rapid
relaxation) would produce similar amount of grain
boundary cracking, resulting in the same reduction of the
creep-fatigue life. Although the main factor in the drop in
the fatigue life when a hold is introduced may be the creep
damage promoting intergranular crack propagation,
oxidation may also play a role.
IV.D. Role of oxidation
One faces a challenge in discussing the role of
environment at elevated temperature. In an ideal case,
comparable creep-fatigue testing performed under an inert
atmosphere would be used as a baseline, exemplifying the
effect of oxidation on fatigue life. However, creep-fatigue
in inert conditions is not practically achievable at high
temperature.
320
a.
b.
c.
Number of cracks
Average length (μm)
Total length (μm)
110
100
90
80
70
60
50
40
35
30
25
20
15
10
5
Proceedings of ICAPP 2011
Nice, France, May 2-5, 2011
Paper 11284
0.3% - air
0.6% - air
0.3% - helium
0
0 500 1000 1500 2000
��������������
Hold Time (sec)
0.3% - air
0.6% - air
0.3% - helium
30
0 500 1000
Hold �������������� time (sec)
1500 2000
2000
1500
1000
500
0.3% - air
0.6% - air
0.3% - helium
0
0 500 1000 1500 2000
��������������
Hold time (sec)
Fig. 10. Grain boundary cracking analysis for LCF and
creep-fatigue testing in air and in helium; a) total number of
cracks, b) crack average length, and c) total length.

Exposure at 950°C in vacuum and in (never
completely) pure helium or argon is accompanied by
decarburization. Dissolution of intergranular carbides to an
increasing depth fosters grain boundary sliding and the
formation of cavities and wedge-cracks in the carbide-free
region 3,18 . These metallurgical instabilities contribute to the
fatigue crack initiation and/or growth, and the fatigue life
may be significantly lower than observed in air 3 . Therefore,
characterization of the cycled alloy microstructure on the
surface and in the surrounding of the cracks will be used to
get an insight of the possible role of oxidation in the fatigue
crack initiation and growth.
Close examination of cross sections of the creepfatigue
tested specimen revealed that surface grain
boundaries were either cracked or showed precipitation of
finger-like alumina (see Figure 7a). Intergranular alumina
precipitates are typical of Alloy 617 oxidized in air or in an
impure but oxidizing helium (as is the process gas in the
present study): the oxygen potential at the scale/alloy
interface, set by the dissociation of chromia, is very low but
high enough for aluminum as well as silicon to be
oxidized 8 . Precipitation is preferred at grain boundaries
which are short diffusion pathways for oxygen. It is
believed that the alumina veins may act as preferential sites
for crack initiation. Once initiated the cracks will open,
exposing new metal surfaces to the gas phase and oxidation
of the crack edge would occur with growth of a chromiumrich
oxide. Eventually, the chromium oxide may fully fill
the crack, forming a large region filled with chromium
oxide.
Following an equivalent process, formation of
intergranular alumina was observed at grain boundaries
intersecting with main oxidized cracks, including ahead of
the crack tip, as shown in Figure 7b. Cracks are likely to
propagate preferentially through these brittle alumina
precipitates and the crack growth may be accelerated.
Therefore, internal oxidation of alumina first below the
external chromia scale and then ahead of the crack tip may
increase both crack initiation and propagation rate. This
hypothesis is consistent with the observed intergranular
cracking and may account for at least a portion of the
reduction in the fatigue life with a tensile hold
IV. CONCLUSIONS
This paper reviews two previous publications by
Carroll and co-workers 10,11 on creep-fatigue testing of Alloy
617 in air and in a VHTR simulated impure helium at
950°C at total strain ranges of 0.3% and 0.6%. This work
supports acceptation of Alloy 617 in the nuclear section
(section III) of the ASME code for application at high
temperature as well as licensing of VHTR systems with
IHX. The creep-fatigue behavior of Alloy 617 at 950°C
may be influenced by both the environment and the creep
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Nice, France, May 2-5, 2011
Paper 11284
deformation that occurs during the tensile hold. The shift in
crack initiation mode from transgranular during fatigue
deformation to intergranular during creep-fatigue
deformation likely results from the environmental
influence. Oxidation may also play a role in determining
the crack propagation mode, tending the creep-fatigue
crack propagation to intergranular as opposed to
transgranular. However, longer hold times did not result in
the further decrease in cycle life suggesting that the
environmental influence on creep-fatigue crack propagation
is not the controlling factor. Measurements of the oxidation
phenomena are planned for a more quantitative approach.
Constant accumulated creep damage despite an increasing
hold time agrees well with the experimental tendency of
cycles to failure saturating. The creep damage in the alloy
bulk, in the form of grain boundary cracking, may have
accelerated the crack propagation. The amount of grain
boundary cracking was similar among all of the hold times
due to the rapid saturation during the stress relaxation.
Creep-fatigue testing will be performed under another
helium environment producing different microstructure and
corrosion products to determine any change in the
deformation mechanisms.
ACKNOWLEDGMENTS
The authors would like to acknowledge Joel Simpson,
DC Haggard, Dave Swank, Randy Lloyd, Tammy
Trowbridge, Todd Morris and Barry Rabin for conducting
the experiments and the microscopy. This work was
supported through the U.S. Department of Energy Nuclear
Energy.
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