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Creep-fatigue interaction of inconel 617 at 950°C in simulated nuclear reactor helium
MaterialsScienceand Engineering, A104 (1988) 37-51
37
Creep-Fatigue Interaction of lnconel 617 at 950 °C in Simulated Nuclear Reactor Helium K. BHANU SANKARA RAO*
Departmentfor ReactorMaterials, NuclearResearch CentreJiilich, Postfach1913, D-5170,Jiilich(F.R.G.) H.-P. MEURER
lnteratom G.rn.b.H., D-5060Bergisch Gladbach 1 (F.R.G.) H. SCHUSTER
Departmentfor ReactorMaterials, NuclearResearch ('entreJiilich, Postfach1913, D-517OJiilich (F.R.G.) (Received May 20, 1987; in revised form February 24, 1988)
Abstract
Strain-controlled fatigue tests have been conducted in impure helium, simulating the primarycircuit coolant of a high temperature gas-cooled reactor to ascertain the influence of strain rate (~ = 4 x 10 --~ to 2 x 10 -5 s- 1), hold condition atpeak strains (tension-only, compression-only and tension-plus-compression holds) and hold time (up to 120 rain) on the low cycle fatigue behaviour of Incone1617. A strain range of 0.6% and a temperature of 950 °C were employed for all the tests. Microstructura! changes which occurred during fatigue deformation were evaluated and damage mechanisms which influence fatigue life identified. A small reduction in fatigue life was found with decreasing L Irrespective of the position of hold at peak strain in a cycle, the hold time always reduced the fatigue life in comparison with continuously cycled tests at lower strain rates but of equal cycle duration. Tensile holds were found to be most damaging, followed by compression holds. Symmetrical tension-plus-compression holds led to fatigue lives which were very close to those of the continuously cycled tests. In the continuously cycled tests with strain rates down to ~ =6.7x 10 -5 S -1, failure was always transgranular with no indication of creep damage. However, for tests with tensile holds, creep damage was evidenced by grain boundary cavitation and oxidation, which were responsible for a reduction in the fatigue life. The formation of thick oxide scales at the surface observed at longer *Permanent address: Materials Development Laboratory, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, Tamil Nadu, India. 0921-5093/88/$3.50
hold times led to chromium-depleted surface zones in which the carbide precipitates were dissolved. The loss of carbides increased grain boundary sliding which enhanced the formation of grain boundary cracks. These cracks shortened the critical length of surface fatigue cracks which, during fracture of the specimen, linked with the intergranular creep cracks in tensile hold tests. The damaging effect of compression hold was attributed to increased inelastic strain and deformation ratcheting found in the cycle; failure occurred by local accumulation of tensile plastic strain which finally caused tensile necking. 1. Introduction This study was performed as part of the materials programme for the development of the high temperature gas-cooled reactor (HTGR) [1, 2]. Inconel 617 is the principal candidate alloy for heat exchanger components, in which low cycle fatigue (LCF) loadings occur owing to reactor start-up and shut-down procedures and to power transients. Operating temperatures are in the range where creep deformation occurs so that the interaction between L C F and creep is an important aspect of the materials behaviour. At the anticipated operating temperatures of an H T G R , only limited data were available on the L C F behaviour of the candidate material Inconel 6 1 7 t [3-9]. Most of these tests were conducted at tThe German standard designation is Werkstoff Nr. 2.4663 or NiCr22Co12Mo; Inconel 617, the designation given to this material by INCO, the company which introduced the alloy,is used throughout this paper. © Elsevier Sequoia/Printed in The Netherlands
38
strain rates of approximately 10-3 s-l, which is orders of magnitude higher than thermally induced strain rates in the heat-exchanging components. There was therefore a need for fatigue data at lower strain rates especially under creep-fatigue loading and for a full understanding of deformation and fracture mechanisms during high temperature fatigue. Therefore, this investigation was carried out to evaluate the effects of strain rate and hold time on the LCF life of Inconel 617 at 950°C in simulated primary-circuit helium gas environment of an HTGR and to correlate the results with the precipitation behaviour, deformation and damage mechanisms in the material, particularly under various hold time loading conditions.
that the temperature profile remained _+3°C throughout the tests, which were conducted in a pre-evacuated chamber and in flowing impure helium gas of composition specified in Table 2. This composition was agreed on by the German HTGR project partners as the representative environment of the primary circuit. A high resolution gas chromatograph was connected to the input and output of the pressure chamber to monitor the gas composition. The test gas was obtained as a bottled premix. The pressure and flow rate of the gas were 1.8 bar and 3 1 h-1 respectively.
2. Experimentalprocedures i
The Inconel 617 test material was supplied by Vereinigte Deutsche Metallwerke AG., Werdohl, ER.G., under the trade name Nicrofer 5220 Co as hot-rolled and solution-annealed plate 17 mm thick, which had been solution treated for 1 h at 1160°C. The chemical composition is given in Table 1 and the grain size was ASTM 6.5. Specimens were machined to the form shown in Fig. 1 and LCF tests were carried out in the laboratories of Interatom G.m.b.H., D-Bensberg, using MTS servohydraulic test machines of + 1 0 0 kN dynamic load capacity. All experiments were performed under total strain control; the axial strain was measured with an extensometer mounted on the ridges of the test specimen. A nominal strain range of 0.6% and a temperature of 950 °C were employed for all the tests. Uniform heating of the specimen gauge length was accomplished using an induction coil and by carefully monitoring and controlling specimen temperature with Pt-(Pt-Rh) thermocouples fixed in a small hole near the gauge section. Previous calibration with dummy specimens showed
TABLE 1
Impurity p (/~bar)
detail X
_____A Fig. 1. Geometry of the specimen (all dimensions are in millimetres).
Chemical composition of Inconel 617
Element Amount(wt.%)
TABLE 2
I
Ni 54.9
Cr 21.44
Co 11.81
Mo 9.34
AI 1.11
Ti 0.55
C 0.087
Partial pressures of the impurities in the helium gas Hz 500 + 50
CH 4 20 _+5
H20 1.5 -+ 1
CO 15 + 5
CO 2 1.0
N2 5
39
Two types of test were conducted. In the tests for examination of the strain rate effect on fatigue life, a symmetrical continuous triangular waveform was used at strain rates of 2 x 10 -5 s ] to 4 × 10 -3 s-1. In the hold time tests for the study of the creep-fatigue-environment interaction the strain was held constant in a cycle at tensile peak strain, at compressive peak strain, or symmetrically at tensile and compressive peak strains. The peak strain holds were superposed on a triangular wave with g = 4 x 10- 3 s- 1. The waveforms employed are schematically described in Fig. 2. Hold time effects were studied for durations ranging from 5 s to 120 min. In order to facilitate the comparison of fatigue properties, strain rates and hold times have been chosen in such a way that for most of the tests the time for 1 cycle remained equal for different types of loading. For example, the cycle durations using 10 min ten-
CONTINUOUSCYCLING I
|
- -
/x,
|
l
TIME
TENSIONHOLD [
-th
t,)
I
TIME
COMPRESSIONHOLD (to)
i /~
I
sion, using 10 min compression, or using 5 min tension plus 5 min compression hold times and the cycle period for the symmetrical triangular waveform tests with g = 2 x 10-5 s 1 are all equal (Fig. 2). Fracture surfaces were examined by optical and scanning electron microscopy (SEM) to determine the crack initiation and propagation modes. Studies were also conducted on longitudinal sections of failed specimens to characterize the behaviour of surface cracks and the microstructural changes that occurred during fatigue testing. These studies were supplemented by transmission electron microscopy (TEM) investigations for which samples were obtained 1 mm below the fracture surfaces normal to the loading axis. Thin foils were prepared by initially grinding the samples down to 250 /~m, followed by electropolishing in a solution containing 20% H 2 S O 4 and 80% methanol at 25 V and 0°C in a twin-jet apparatus.
/
TIME
ENSION AND COMPRESSIONHOLD ~-- th-~
"Z Z
E
°
/
TIME
] - q c --'~
--cycle period-Fig. 2. Wave shapes employed in LCF testing.
3. Results
3.1. Determination of the number N l of the cycles to macrocrack initiation and number N v of cycles to failure N L and N v were estimated from the plots of absolute peak stress ratio IO'compression/O~ensionl IS. number of cycles. The number of cycles corresponding to a 2% and 50% increase from the plateau value of the stress ratio curve are taken as the life NI to macrocrack initiation and the life Nv to failure respectively. However, the determination of N~ and NF could not be established by this procedure under compression hold conditions because no variation in stress ratio occurred up to the end of specimen life owing to identical stress responses in the tension and compression portions. In this particular case, Nv is taken as the life at rapid onset of tensile stress drop that took place towards the end of the test. 3.2. Influence of test variables on fatigue life A summary of the experimental results showing the influence of strain rate, different types of hold time and different durations of hold time on NI, Nr and time te to failure is presented in Table 3. Figure 3 shows the relationship between the fatigue life reduction factor R and the cycle period for different loading conditions. The fatigue life reduction factor is defined as the ratio of Nv/Nvo, where N v is the life recorded for a
40 TABLE 3
Influence of strain rate, hold condition and hold time on fatigue life at
Specimen number
g in ramp (s- 1)
1694 1437 1724 1727 1729 1454 2131 1699 1450 1712 1449 1698 1445 1713 1707 1708 1716 1715 1717
4.0 × 1 0 - 3 4.0 × 1 0 - 3 4.0 × 1 0 - 4 2.0 × 10 -+ 6.7 x 10 -5 2.0 x 10 -5 4.0 x 10 -3 4.0x 10 3 4.0 x 10 -3 4.0 x 10 -3 4.0 X 1 0 - 3 4.0 x 10 -3 4.0 x 10 -3 4.0 x 10 -3 4.0x 10 3 4.0x 10 3 4.0x 10 3 4.0x 10 3 4 . 0 x l O -3
l
i0 °
'
,
.,f
HoM condition
q
Tension Tension Tension Tension Tension Tension Tension Tension Tension Compression Compression Tension + compression Tension + compression
,
i
, I
w
'
' ' 1
'
,
I,
O tension hold times 13 compression held times
--0
__~__~
..
"n"*,7:22;
V etrl|n
"'"
10 -I
10-1
Fig. 3. The life reduction factor R against the cycle period in minutes for different loading conditions. R is defined as the ratio of N~-/Nvo where N F is the fatigue life recorded for a given strain rate g or wave shape and hold time and Nv0 is the reference life value for continuous cycling with g = 4 × 10-3 s-I (o).
g i v e n s t r a i n r a t e o r h o l d t i m e t e s t a n d Nv0 is t h e r e f e r e n c e life v a l u e f o r c o n t i n u o u s c y c l i n g at g=4×10-3s -1. The results of continuous-cycling tests showed a v e r y g r a d u a l r e d u c t i o n in f a t i g u e life w i t h t h e l o w e r i n g o f t h e i m p o s e d s t r a i n r a t e , i.e. w i t h a n i n c r e a s e in c y c l e d u r a t i o n . A r e d u c t i o n in g b y a f a c t o r o f 2 0 0 c a u s e d a f a t i g u e life r e d u c t i o n o f less t h a n a f a c t o r o f 2.
T =
9 5 0 °C (total strain e t = 0.6%)
HoM time (rain)
Number N I of cycles to crack initiation
Number N~ of cycles to failure
Time to failure (h)
0 0 0 0 0 0 0.083 1.0 1.0 3.0 3.0 10.0 10.0 30.0 120 1.0 10.0 0.5 +0,5 5.0 + 5,0
2530 3000 1900 1610 1775 1250 1080 610 700 335 560 270 300 220 130 -----
3025 3200 2725 2370 2325 1580 1420 830 890 650 740 450 465 380 215 1030 665 2200 1500
2.5 2.7 22.7 39.5 116.3 263.5 3.5 14.5 15.0 32.5 37.0 15.5 77.5 190.0 430.0 17.5 111.5 367.0 251.5
Hold times at the tensile peak strain led to a pronounced reduction in fatigue life. It can be seen that tensile hold periods of only 5 s resulted in shorter fatigue lives. A compressive peak strain hold was also found to cause large reductions in fatigue life but its effect was slightly less than that of tensile hold. Equal hold periods both in tension and in compression yielded reduction factors which were very close to the continuouscycling data with identical cycle duration. It should be noted that fatigue life reduction increased continuously with an increase in hold time irrespective of the type of hold conditions employed in a given cycle; no saturation of the reduction was evident. 3.3. Cyclic stress response as a function of strain rate and type of hold time Figure 4 shows half-fife hysteresis loops for various testing conditions to characterize and compare their effects on material cyclic deformation behaviour. Continuous strain cycling promoted approximately symmetrical stress response in the tension and compression portions of the cycles (Fig. 4(a)). Reducing the strain rate yielded a considerable decrease in the maximum tensile and compressive stress and increased the amount of inelastic deformation in each cycle. Irrespective of the strain rate applied, the stress approached a constant level. This feature has also been found in
41
tensile deformation behaviour of Inconel 617 and other high temperature alloys at 950 °C [10, 11]. At the strain reversals during the continuouscycling experiment with a strain rate of 4 × 10 -3 s -1 (Fig. 4(a) and Table 4) a rapid stress relaxation was observed, indicating an inelastic strain rate faster than the applied total strain rate. At low strain rates, i.e. for g~< 2 × 10 -4 s -~, this rapid relaxation effect was absent because of lower initial stresses. 250 20O
f
f
t 50 I00
~.
50 o
"~ -5o - ~:2xlO-Ss ~
-I00 -150
-200
: 2 x 10-4 s-I
,
-250 (a)
0.1%
250 200 150 I00
o2
5o o
b
-50 -I00 -150 -200 -250 (b)
continuous cycling
IO rnln tenslon hold
5ram tension + 5 rnln compression hole
s "--~
~ _ 0.1%
Fig. 4. Stress-strain hysteresis loops (a) representing the midlife to crack initiation in continuous-cycling tests at different strain rates and (b) taken at mid-life to crack initiation for the continuous-cycling condition ( g = 4 x 10 _4 s - l ) , 10 rain tension hold, 10 rain c o m p r e s s i o n hold and 5 min symmetrical hold.
When hold times were applied, in either tension or compression, a very pronounced stress relaxation occurred (Fig. 4(b)). Even though only short hold times have been applied, the stresses relaxed to nearly zero. In the cycles consisting of tension-only or compression-only hold times, this pronounced stress relaxation yielded asymmetrical hysteresis loops. For tension hold times the inelastic strain in the compressive part of the hysteresis loop was enlarged compared with the inelastic strain in the tensile part and the opposite correlation was found for compression hold times. The cyclic stress response curves for various testing conditions are depicted in Fig. 5. In this figure the tensile stress amplitude is plotted against the fraction N/Nv of fatigue life in order to show more clearly the various deformation stages. The onset of the stress drop towards the end of the test generally indicates the formation of macrocracks and their subsequent growth [12]. Continuous-cycling tests exhibited, after a rapid initial stabilization of peak stresses, a more or less steady state stress response over a large portion of fatigue life (until N / N F <<.0.70). Tests carried out with a 10 min tension hold showed a small rapid initial softening followed by a gradual softening until macroscopic crack growth occurred. It should be pointed out that the relative fraction of fatigue life spent in lhe crack propagation stage was considerably longer in the 10 min tensile hold condition than in the continuous-cycling test with g = 4 x 10 3 s- l Compression holds alone and tension-pluscompression holds initially caused slight cyclic strain hardening followed by softening. Symmetrical holds developed slightly higher tensile peak stress than the holds introduced in compression or tension alone.
TABLE 4 Computed values of elastic strain Aeel , plastic strain Aep, rapid relaxation strain epr, and effective plastic strain Aep' as a function of strain rate
Specimen number 1694 1724 1727 1729 1454
g (s 1)
4x10 3 4 x 10 -~ 2 x 10 4 6.7 x 10 5 2.0 x l O- 5
Aet (%)
0.6 0.6 0.6 0.6 0.6
Aeel (%)
0.21 0.17 0.16 0.13 0.09
Ae (%3
0.33 0.39 0.43 0.47 (I.49
Rapid relaxation strain 6pr (%)
Effective plastic strain
In tension
In compression
Aep' = Aep + 2e F
0.03 0.015 ----
0.03 0.015 ----
0.39 0.42 0.43 0.47 0.49
Ae~ is the total strain range, A e p = A e , - 2 d I E and %r = -IAddl/E w h e r e E is Young's m o d u l u s , d the stress amplitude in the cycle and A d d the stress drop at strain reversal.
42
• 4. I0 3 s 4
1
~lOO 0 tOmin tension hold
200
leo
O 200
I0 rain compressionhold.
;,--____._.._
100 0 200
5min tension +
leo
0 200
Fig. 6. Photograph of broken specimens tested at g= 4 × 10-3 s t under the followingconditions: (a) continuous cycling;(b) 10 min tension hold; (c) 10 min compression hold; (d) 10 min tension-plus-compressionhold. Geometrical instability in the gauge portion of the samples subjected to tension-onlyand compression-onlyholds can be clearlyseen.
k
~= Z' 10.4 S"1 10l O 20| E=2.1OSs "1
Alloy 800H in hold time tests conducted at
I0|
850 °c [8]. 0
0,1
i
i
i
I
,
i
i
,
0,2
0,3
0.4
0.5
0.6
0.7
0.8
0.9
- -
I,O
N/NF
Fig. 5. Cyclic tensile stress response curves for various loading conditions studied. The tensile stress amplitude is plotted against the fraction N/Nv of fatigue life. The total strain range Ae is 0.6% throughout. By decreasing the strain rate from 4 x 10-3 to 2 x 10 -5 s -], the value of the saturation stress was reduced from 228 to 64 MPa. Such a strong g dependence of saturation stress was also noted in the LCF tests conducted in air at 850°C [13]. In hot tensile tests with Inconel 617 at 950 °C the ultimate tensile strength shows the same dependence on strain rate as the saturation stress during the L C F tests [11, 14]. 3.4. Fractographic observations The effect of loading conditions on the physical appearance of the failed specimens is shown in Fig. 6. There was no significant departure from the original geometry after L C F testing with continuous-cycling waveform (Fig. 6(a)) or symmetrical hold (Fig. 6(d)). It can be seen that tensile hold (Fig. 6(b)) induced geometrical instability into the gauge length in the form of bulging whereas in compression hold (Fig. 6(c)) the instability was manifested in the form of necking. Similar observations have been made on Inconel
In the continuous-cycling tests a clear transition in fatigue crack initiation behaviour was found from transgranular stage I shear cracking at g > 2 x 10- 4 s- t to intergranular initiation at g < 6 . 7 x 10 -5 s -1. Crack propagation was found to be transgranular stage II at all the strain rates in continuous cycling. In the 1 min tension hold test, cracks initiated transgranularly by stage I shear cracking on the specimen surface (Fig. 7(a)) whereas crack propagation occurred by mixed mode (Fig. 7(b)). In the test with tensile holds of greater than 10 min, crack initiation occurred intergranularly (Fig. 7(c)). The fracture surface of the samples with tensile hold periods of longer than 1 rain contained a fairly large number of grain boundary cracks in the interior of the specimens. The compression hold test specimens displayed a dimple fracture (Fig. 7(d)) similar to a tensile fracture accompanied by necking. From the small remaining fracture surface in compression hold tests the crack initiation mechanism could not be deduced. The fracture surface of symmetrical hold tests showed transgranular stage I crack initiation and mixed-mode propagation. 3.5. Precipitation behaviour Microstructural features such as carbide precipitation can play an important role in cyclic
43 stress-strain behaviour, the modes of crack initiation and propagation, and fatigue ductility. In the present study, this aspect was investigated using electron diffraction, energy-dispersive X-ray analysis (EDXA) and SEM backscattered electron image techniques.
3.5.1. Primary carbides Two types of carbide not dissolved by the solution treatment were found in all the samples studied. Relatively coarse cube-shaped particles were identified as primary Ti(C, N) particles by EDXA. The second type of particle appeared as stringers in both intragranular and intergranular regions (Fig. 8). An E D X A of these particles indicated that they contained both chromium-rich and molybdenum-rich regions corresponding to M 2 3 C 6 and M 6 C types of carbide, which have been identified as the grey and white areas respectively. The proximity of these phases confirms that M6C is a result of M23C~ transformation [15, 16].
Fig. 7. (a) Fracture surface of a 1 min tensile hold sample, transgranular stage I crack initiation and stage II propagation; (b) mixed-mode propagation in the interior of a 1 min tensile hold specimen; (c) intergranular crack initiation in a 120 min tensile hold sample; (d) fracture surface showing dimple fracture in a 10 min compression hold test.
3.5.2. Development of secondary carbides as a function of loading condition The dark field TEM micrograph shown in Fig. 9(a) illustrates the morphology and distribution of M23C6 precipitated along the grain boundary in the samples tested at a strain rate g of 4 x 10 -3 s-1 with zero hold. These carbides had an f.c.c, ordered crystal structure coherent with one of the neighbouring grains, as shown by electron diffraction. One of the very fine intergranular M23C6 particles is shown in Fig. 9(b). No secondary M~C was found in samples continuously cycled at g= 4 x 10 -3 s-l; this is consistent with the prediction of an incubation time for its appearance at 950°C [16]. The specimen life t~of 2.7 h for g= 4 x 10 -3 s- ~is far less than the incubation time reported [16] for onset of M~C precipitation. At g < 4.0 x 10 - 4 s- ~, copigus amounts of M~C precipitated together with M23C6 were found in both intragranular and intergranular regions (Fig. 9(c)). The introduction of a 1 min tensile hold at the peak strain led to an increase in the volume fraction and size of intragranular M23C6 (Fig. 10(a)). An increase in both the size and the density of grain boundary carbides was also noted. The carbides formed along the grain boundary were predominantly M23C6, a small amount of M 6 C w a s also detected. In addition to the grain boundaries,
44
~
! 20 Dim
II
Fig. 8. SEM image obtained in backscattered electron mode showing undissolved massive precipitates in intragranular regions. White precipitates correspond to MaC and grey precipitates to M23C 6 (T=950°C; ramp rate ~=4x 10 3 s-~; tensile hold for 10 min). deformation bands and twin boundaries were also preferential sites for carbide precipitation (Fig. 10(b)). Secondary precipitation of finely dispersed Ti(C, N) was detected in tests with tensile holds of 1 rain or longer in the intragranular regions. A further increase in tensile hold to 10 min resulted in the rapid coarsening of grain broundary precipitates and increased interparticle spacings (Fig. 10(c)). Raising the duration of hold to 120 min did not produce any further significant changes in the grain boundary precipitate size, density, morphology and interparticle spacings. The SEM backscattered electron image in Fig. 1 l(a) illustrates the typical microstructure obtained in samples tested with a 5 min tension plus 5 min compression hold. It can be seen that discrete and large precipitates of M6C (white phase) and M23C 6 (grey phase) coexisted on large-angle grain boundaries, but few were found in the matrix. The proportion of M6C on grain boundaries in the 5 min symmetrical hold test was very much higher than that observed in either the longest-duration tensile hold or the 10 min compression hold test.
Fig. 9. (a) Dark field TEM m]crograph illustrating grain
boundary M23C6 (T=950°C; g=4xl0 -3 s-I); (b) intra-
granular M23C6 carbides with interracial dislocations (T= 950°C; g= 4 × 10 -3 S-I); (C) backscattered SEM image showing Ti(C, N) (dark phase), M23C6 (greyphase) and M6C (white phase) particles (T= 950°C; g= 2 x 10 -5 s- i). In the 10 min compression hold test a nearly continuous network of grain boundary carbides was observed, as shown in Fig. 1 l(b). In the tests with holds of longer duration, intragranular M23C6 and M6C acquired a spherical shape (Fig. 1 l(c)) compared with the cuboidal shape which was commonly observed in continuous-cycling tests or tests with a short hold time. The formation of )/precipitates (Ni3(A1, Ti)) in Inconel 617 has been described in the literature [16, 17]. In the present study, the introduction of
45
~f
E ¸¸¸
Fig. 10. (a) TEM micrograph showing intragranular M23C6 particles in the samples tested with a 1 min tension hold; (b) secondary-electron SEM image showing precipitation on deformation bands in the samples tested with a 1 min tension hold; (c) intragranular and intergranular precipitate distribution in samples tested with a 10 rain tension hold.
a compression hold alone or in combination with a tensile hold induced serpentine-shaped ~' phase identified by diffraction, but the amount detected was so small that no effect on ductility and fatigue life is expected.
3.6. Deformation behaviour Regardless of the loading condition, the alloy exhibited generally subgrain formation in the
Fig. 11. (a) SEM backscattered electron image illustrating M23C6 (grey phase) and M6C (white phase) precipitates in the samples tested with a 5 min symmetrical hold; (b) secondaryelectron SEM image showing continuous network of grain boundary carbides in the samples tested with a 10 min compression hold; (c) bright field TEM micrograph of spherical M23C6 precipitates in the samples subjected to a l0 min compression on hold (the precipitates are found to be spherical at longer hold times in contrast with cube-shaped particles seen at shorter holds).
majority of the grains (Fig. 12(a)). Multiple slip in some of the grains was also observed, suggesting that the mode of deformation was predominantly homogeneous at 950 °C. In the tests with a tensile hold, the alloy also showed some tendency to form dislocation tangles in the matrix and dislocation pile-ups at the grain boundaries (Fig. 12(b)). However, in spite of dislocation pile-ups, no intergranular cracking could be observed. Under a 5 min symmetrical hold, the substructure
46
Fig. 13. Transgranular secondary crack in the specimen tested with g= 4 x 1 0 - 3 s - 1 . The arrow in the top right-hand corner indicates the loading direction. 4. Discussion
4.1. Influence o f the strain rate on fatigue life reduction
Fig. 12. Bright field TEM micrographs showing(a) substructure in samples tested with g= 4 x 10 3 s- ', (b) the pile-up of dislocations on grain boundaries and dislocation tangles in the intragranular regions and (c) interaction between coherent M23C 6 and dislocations (carbide precipitates have been identifiedthrough moire fringe contrast). indicated extensive interactions between coherent precipitates and dislocations (Fig. 12(c)). In this figure the coherent M23C6 precipitates were identified by moire fringe contrast. It is reasonable to correlate such interactions with the slight cyclic hardening observed for this loading condition.
It is evident from Fig. 3 that decreasing the strain rate results in a small reduction in fatigue life. Associated with this reduction are a decrease in stress amplitude, an increase in inelastic strain range (i.e. an increase in the width of the hysteresis loop at zero stress; see Fig. 4 and Table 4) and a change in crack initiation mode from transgranular at g >t 2 x 10-4 s- 1 to intergranular at g ~ < 6 . 7 x l 0 - S s -1. At the lowest strain rate, cracks were initiated by the oxidation of surfaceconnected grain boundaries. Intergranular cracks propagated to only one grain diameter before a transition to transgranular propagation took place. From the studies of secondary cracks in longitudinal sections, it was found that the transgranular crack propagation path remained sharp and narrow (Fig. 13). This observation suggests very little interference of the oxidation process in the crack propagation mechanism. Furthermore, microstructural investigations revealed no evidence of any creep damage in the bulk even at the lowest strain rate investigated in continuouscycling tests. On the basis of detailed investigations of the influence of strain rate at 850 °C in air, it has been argued that the reduction in fatigue life with decreasing strain rate results from the combined effects of increased inelastic strain and from oxidation of surface connected grain boundaries and environmentally accelerated mixed-mode propagation [13]. In the range of strain rates examined in this investigation, the contribution from environmental attack to reduction in life appeared only at g < 6 . 7 x 10 -5 s -1 in the form of oxida-
47
tion-assisted intergranular crack initiation. Moreover, the reduction in fatigue life occurred with decreasing g, even in the range of strain rates (~= 2 x 10-4-4 x 10 -3 s -1) where entirely transgranular crack initiation and growth prevailed. These observations demonstrate that fatigue life is largely governed by the increase in effective plastic strain component of the cycle as g is lowered. 4.2. Hold time effects 4.2.1. Mean stress under hold time conditions From Fig. 4(b), it can be observed that the addition of either tension-only or compressiononly holds causes no significant mean stress in the hysteresis loops. The capacity of the material to withstand the mean stress depends on the ratio of inelastic to elastic strains in a cycle [18]. If the inelastic strain is large relative to the elastic strain, as in the case of the present experiments, the mean stress which can develop is very low and consequently mean stresses do not influence cyclic life significantly. 4.2.2. Relaxation The damage behaviour during the hold period could be best described by characterizing the stress relaxation behaviour as a function of time. A typical mid-life stress relaxation response curve for the 120 min tensile hold test is shown in Fig. 14. It can be seen that rapid relaxation down to half of the maximum stress occurred in less than 1 s. Further, the inelastic strain rate gr associated with relaxation strain continuously decreased with increasing hold time from approximately 6.5 x 10 4 t o 4 × 10 9 s - l .
In Fig. 15 the inelastic strain rates associated with stress relaxation for different hold conditions were compared with minimum-creep-rate data available from creep tests for various stress levels on the same heat. The strain rates during stress relaxation are generally higher than the corresponding minimum creep rates at all the stress levels. The shape of relaxation rate curves resemble the shape of the minimum creep rate vs. stress curves at a lower strain rate but distinctly deviate at higher rates. Inelastic strain rates found in the rapid relaxation period correspond to those which are expected for precipitation-free matrix deformation, while those observed in slow relaxation period are typical for creep deformation. 4.2.3. Cavitation damage in tensile hold tests The build-up of tensile inelastic relaxation strain caused severe cavitation damage within the bulk of the material. Figure 16(a) clearly illustrates the development of intergranular cavities (R type) associated with grain boundary precipitates adjacent to the fracture surface in a 1 min tension hold test. Figures 16(b) and 16(c) describe the cavity formation and cavity linkage on grain boundaries in the samples cycled with 10 and 1 20 rain tension holds in the regions remote from the fracture surface. R-type cavities form at grain boundaries and second-phase particles.
n0, s0 100 lu e00 ~s0 ~ ~ ,~ ' , ' ~ , ' i " ~ ",0~,~',',?' /~'//~-'~--"~i ~ v~ llimin l~.~n.~u |~ tension
!
////°
,.o
........////j
. .........
............. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Ill
....
;,,,j,
....
no
0
o lo'
ao'
m'
lo't/sec~
~o'
Fig. 14. Mid-life stress relaxation curve in a 120 rain tensile hold test (it, inelastic relaxation strain rate),
~ .... 50
i .... IH
I .... 150
i
. . . .
ZooG./Mra25o
Fig. 15. Relaxation strain rate data for different hold conditions. M i n i m u m - c r e e p - r a t e data from stress rupture tests are also included in the figure for c o m p a r i s o n wilh relaxation strain rate data.
48
Fig. 17. (a) Grain boundary cracks originating from wedge cracks in the near-surface regions free from carbide precipitation (10 min tension hold); (b) thick oxide layer formed on the surface (the crack faces of secondary cracks are also oxidized, and internal oxidation ahead of the crack tip can also be seen) ( 120 min tensile hold).
Fig. 16. (a) Intergranular cavitation damage adjacent to the fracture surface in the sample tested with a 1 min tension hold (cavities are associated with grain boundary precipitates); (b) cavity coalescence on grain boundaries oriented at 90 ° to the loading axis (10 min tension hold); (c) intergranular cavity damage in the samples subjected to a 120 min tension hold.
In addition to the bulk cavity damage, a large number of grain boundary cracks were observed in the surface regions particularly when the tension hold times were more than 10 min (Fig.
17(a)). A closer examination showed that these cracks formed in the near-surface regions which were free from M23C6 precipitation nucleate as wedge-type cracks. This observation supports the conclusion that sliding is higher at grain boundaries free from carbides [19] which facilitates linking of cavities to form grain boundary cracks. Furthermore, the damage due to grain boundary oxidation was found to be considerable at longer tensile holds. Figure 17(b) shows the surface oxide layer and oxide-filled subcracks in the failed sample with an imposed tensile strain hold of 120 min. A n E D X A of the oxide layers demonstrated that these were basically of chromia. From Fig. 17(b), it is also apparent that the oxygen diffusion ahead of the crack caused internal oxidation of aluminium. A zone of 10/~m from the surface was also found to contain discrete particles of aluminium oxide. Since most of the oxides are intrinsically brittle, the fatigue
49
cracks in tests with longer hold periods initiated in oxides which are formed preferentially at surface-connected grain boundaries. The oxidation-induced surface intergranular cracks penetrated deeply into the interior and merged with independently formed intergranular wedge cracks in the near-surface regions and R-type cavities in the bulk. 4.2.4. Compression hold tests Compression hold tests did not exhibit creep cavities in the specimen interior. In addition, none of these features such as a thick surface oxide layer, carbide-free zones and wedge cracking in the near-surface zones which have been prominently displayed in long period tensile hold tests was found in compression hold tests (Fig. 18). Bulk creep damage is unlikely to develop during the periods of compressive stress relaxation. Furthermore, compressive holds have been generally acknowledged as capable of sintering creep cavitation damage. Therefore the damaging effect of compression hold on life reduction could only be ascribed to the increased amount of inelastic strain in a cycle resulting mainly from stress relaxation. It must be remembered that, during compression holds, test specimens exhibited geometrical instability (Fig. 6) and produced tensile necking failure. This has occurred in the absence of significant tensile mean stress. The reason why compression hold strain cycling leads to tensile fracture has been interpreted by Manson [20]. The progressive accumulation of tensile inelastic strain with continued cycling in the tensile half and its detrimental effect compared with the accumulation of relaxation creep strain during compression hold leads to tensile fracture.
4.2.5. Symmetrical hold tests The optical micrographs presented in Fig. 19 illustrate the results for 5 min symmetrical hold tests. In contrast with the uniformly distributed intergranular cavitation damage in the specimen interiors of tensile hold tests, the cavity damage has been confined close to the fracture surface (Fig. 19(a)). No wedge cracks were seen in the carbide-depleted zone in the near-surface regions (Fig. 19(b)) as opposed to the observations in tensile hold tests. The few intergranular cracks which have merged with surface cracks appeared to be a consequence of internal oxidation of the grain boundaries. The non-occurrence of wedge cracking in the carbide-depleted region would increase the critical length of the fatigue crack before the onset of interaction with the cavity damage in the central region of the specimen is established. This interaction and the increased amount of inelastic strain in the cycle are responsible for the slightly inferior fatigue resistance of the symmetrical hold tests compared with continuous-cycling no-hold tests of equal cycle duration. In spite of the higher amount of inelastic
'o.T: /~
....
•
~. . . . . . . . ,
-
-".~,
~
.'" : "~'" , , ' ~ . 1 / ~ ' ; )
.~-
" ,:
:
" "
~ 1 5 0 p m
Fig. 18. View of the surface and subsurface regions in the longitudinal section of the samples tested with a 10 rain compression hold.
. . . . . . . . . . .
'
'
.................
:
150
Fig. 19. (a) Ca~tation damage near the fracture surface in the 5 min symmetrical hold test; (b) carbide-depleted zone in the surface regions and secondary cracks in this zone.
50 strain in a cycle and fatigue-cavity damage interaction, symmetrical holds exhibited better fatigue resistance than compression-only holds probably because of the non-existence of deformation ratcheting in the cycle. It has been suggested that, if compressive hold periods follow the tensile hold periods, sintering of cavities formed during the tensile hold period takes place during the compressive hold [21[. In fact, the experimental findings of Majumdar and Maiya [22] on AISI 304 stainless steel showed that symmetrical holds produce only transgranular failure contrary to our observations of an interaction between the fatigue and cavity damage. It should be noted from Fig. 19(b) that the cavities became rather large, corresponding to the size of the grains, and were definitely not associated with grain boundary carbides as in the case of tensile-only holds. Furthermore they were oriented at about 45 ° to the stress axis. These observations indicate that these cavities nucleate and grow by irreversible shear deformation.
4.3. Corrosion effects In this investigation the main corrosion effects observed in simulated primary-circuit helium of the HTGR (Table 2) were the formation of predominantly chromium-based oxide scales at the surface, internal oxidation of aluminium beneath the surface oxide layer, in the adjoining areas of crack faces and at the crack tip, and the carbidefree zone in the near-surface region. These effects are consistent with the data reported on isothermal corrosion behaviour of this alloy at 950 °C in an HTGR helium environment [23, 24]. In the simulated HTGR helium environment (Pra~o above 0.5 pbar) the aluminium content of Inconel 617 is too low for a surface scale based on AIzO 3 to form; A1203 exists only as internal oxide beneath the stable chromium oxide layer [24]. The carbide-free zone in the near-surface region results partially from the depletion of chromium which is consumed in oxide formation. An E D X A of the carbide-free area has shown a chromium-depleted zone of only 30 pm, which is in agreement with the data reported on isothermal exposures in HTGR helium at 950 °C [24]. Because the carbide-free area is much larger than the chromium-depleted zone, it is not reasonable to assume that the depletion of chromium is entirely responsible for the development of carbide-free zones. This indicates that decarburization was occurring in the near-surface regions. In
fact, the analysis of the test atmosphere of the outlet of the LCF test chamber by gas chromatography indicated a considerable increase in CO content in the longer-hold-time tests. Decarburization has also been noted for this alloy after stress-free exposure in HTGR helium atmospheres at 950 °C [23]. 5. Conclusions
(1) Decreasing strain rate caused a reduction in the number of cycles to fatigue crack initiation and to failure. In the range of strain rates examined ( g = 4 x 10-3-2 x 10 -5 s-l), the major factor responsible for fatigue life reduction was the continuous increase in inelastic strain component in a cycle as the strain rate was lowered. The contribution from environmental attack was found only at g<6.7 x 10 -5 s -1 in the form of oxidation-assisted intergrannlar crack initiation. (2) Irrespective of the position of hold at peak strain in a cycle, hold time always reduced fatigue life compared with continuous-cycling tests. (3) Hold periods in tension led to significant reductions in fatigue life; these reductions were slightly larger than those observed for compression-only hold tests and considerably larger than those involving symmetrical hold periods in both tension and compression. (4) Hold periods in compression produced failure by necking. Owing to the large amount of relaxation, an unbalanced tensile inelastic strain occurred and was accumulated during cyclic loading. No cavity damage was observed. (5) The reduction in fatigue life which occurred when tensile hold periods were applied was principally due to the interaction between the surface-initiated fatigue crack and interior grain boundary creep cavitation damage formed during periods of tensile stress relaxation. (6) At longer tensile holds, a large number of grain boundary cracks were found in the nearsurface regions depleted of grain boundary carbides. Cracking in these regions is believed to shorten the critical length of surface fatigue cracks which establishes the interaction with bulk cavity damage associated with the carbides. (7) Oxidation played a dominant role in fatigue initiation and early propagation at longer tensile holds. There was a progressive change in crack initiation mode from transgranular stage I at short tensile hold periods to intergranular
51
initiation at tensile hold periods longer than 10 min. (8) The corrosion effects observed in the HTGR helium atmosphere consisted of the formation of predominantly chromium-based oxide scales at the surface, internal oxidation of the aluminium beneath the surface oxide layer, in the adjoining areas of crack faces and ahead of the crack tip, and formation of carbide-free zones in the near-surface regions. The carbide-free zones formed as a result of consumption of chromium by oxide layer formation and because of decarburization occurring in the HTGR helium atmosphere. Acknowledgments The authors are grateful for the experimental work done by the staff of the Fatigue Laboratory of Interatom G.m.b.H. and the support by the Metallography and TEM Laboratories at the Department for Reactor Materials, Nuclear Research Centre, Jiilich. The careful reading of the manuscript by Mr. E J. Ennis is gratefully acknowledged. K. B. S. Rao is grateful to Shri C. V. Sundaram, Dr. Placid Rodriguez and Dr. S. L. Mannan, all of Indira Gandhi Centre for Atomic Research, Kalpakkam, India, for support during the period of his stay at Jfilich under the Indo-German bilateral agreement. The work has been performed within the framework of the German HTGR project Prototype Plant for Nuclear Process Heat and has been sponsored by the State of Northrhine Westfalia and the Minister for Research and Technology of the F.R.G. References 1 H. Nickel, F. Schubert and H. Schuster, Proc. Conf. on Gas-cooled Reactors Today, Bristol, September 10-24, 1982, BNES, London, pp. 173-178. 2 Status of metallic materials development for application in advanced high-temperature gas-cooled reactors, Nucl. Technol., 66(1984) 11-721. 3 H. E Meurer, H. Breitling and E. D. Grosset, Proc. Conf. on the Behaviour of High Temperature Alloys in Aggressive Environments, Petten, October 15-18, 1979, Metals Society, London, i980, pp. 1005-1016. 4 H. E Meurer, H. Breitling and E. D. Grosser, Proc.
Specialists' Meet. on High Temperature Metallic Materials for Application in Gas Cooled Reactors, in IAEA Rep. 158, May 1981, pp. U 1 - U l l (International Atomic Energy Agency). 5 J.P. Strizak, C. R. Brinkman and E L. Rittenhouse, Proc. Specialists" Meet. on High Temperature Metallic Materials for Application in Gas Cooled Reactors, in IAEA Rep. 158, May 1981, pp. T1-T15 (International Atomic Energy Agency). 6 H. Hattori, M. Kitagawa and A. Ohtomo, J. Soc. Mater. Sci., Tokyo, 32(1983)667-671. 7 M. A. Burke and C. G. Beck, Metall. Trans. A, 15 (1984) 661-670. 8 H. E Meurer, G. K. H. Gnirss, W. Mergler, G. Raule, H. Schuster and G. Ullrich, Nucl. Technol., 66 (1984) 315-323. 9 K. Bhanu Sankara Rao, H. Schiffers and H. Schuster, P~oc. Int. Conf. on High Temperature Alloys -- Their Exploitable Potential, Petten, October 15-17, 198.5, Elsevier, Amsterdam, 1987, pp. 411-422. 10 U. Bruch, E J. Ennis, E. teHeesen, Nucl. Technol., 66 (1984) 357-362. 11 A. Abdel Azim, Jiil-Rep., to be published (Nuclear Research Centre JiJlich). 12 K. Bhanu Sankara Rao, M. Valsan, R. Sandhya, S. K. Ray, S. L. Mannan and E Rodriguez, Int. J. Fatigue, 7(1985) 141-147. 13 K. Bhanu Sankara Rao, H. Schiffers, H. Schuster and H. Nickel, Metall. Trans. A, 19 (1988) 359-371. 14 H. P. Meurer, Interatom G.m.b.H., unpublished research work, 1984. 15 R. F. Decker and C. T. Sims, in C. T. Sims and W. C. Hagel (eds.), The Superalloys, Wiley, New York, 1972, pp. 33-77. 16 H. Kirchh6fer, J. Rottmann, E Schubert and H. Nickel, Jiil-Rep. 1903, March 1984 (Nuclear Research Centre Jiilich). 17 P. P. Schepp, Doctoral Thesis, Universit~t D-ErlangenNiirnberg, 1983. 18 D. C. Lord and U F. Coffin, Jr., Metall. Trans., 4 (1973) 1647-1653. 19 R. yon der Gracht, P. J. Ennis, A. Czyrska-Filemonowicz and H. Schuster, Proc. Int. Conf. on Creep, Tokyo, April 1986, Japan Society of Mechanical Engineers, Tokyo, 1986, pp. 123-128. 20 S. Manson, in A. E. Carder, A. J. McEvily and C. H. Wells (eds.), Fatigue of Elevated Temperatures, ASTM Spec. Tech. Publ. 520, 1973, pp. 744-782. 21 J. B. Conway, R. H. Stenz and J. T. Berling, Fatigue, tensile and relaxation behaviour of strainless steels, Publ. TID 26135, 1975 (National Technical Information Centre, U.S. Department of Commerce). 22 S. Majumdar and P. S. Maiya, Proc. 2nd Int. Conf. on the Mechanical Behaviour of Materials, Boston, MA, August 16-22, 1976, American Society for Metals, Metals Park, OH, 1976, pp. 921-928. 23 W.J. Quadakkers, H. Schuster, Werkst. Korros., 36 ( 1985) 141-150. 24 W.J. Quadakkers, Werkst. Korros., 36(1985) 335-347.
37
Creep-Fatigue Interaction of lnconel 617 at 950 °C in Simulated Nuclear Reactor Helium K. BHANU SANKARA RAO*
Departmentfor ReactorMaterials, NuclearResearch CentreJiilich, Postfach1913, D-5170,Jiilich(F.R.G.) H.-P. MEURER
lnteratom G.rn.b.H., D-5060Bergisch Gladbach 1 (F.R.G.) H. SCHUSTER
Departmentfor ReactorMaterials, NuclearResearch ('entreJiilich, Postfach1913, D-517OJiilich (F.R.G.) (Received May 20, 1987; in revised form February 24, 1988)
Abstract
Strain-controlled fatigue tests have been conducted in impure helium, simulating the primarycircuit coolant of a high temperature gas-cooled reactor to ascertain the influence of strain rate (~ = 4 x 10 --~ to 2 x 10 -5 s- 1), hold condition atpeak strains (tension-only, compression-only and tension-plus-compression holds) and hold time (up to 120 rain) on the low cycle fatigue behaviour of Incone1617. A strain range of 0.6% and a temperature of 950 °C were employed for all the tests. Microstructura! changes which occurred during fatigue deformation were evaluated and damage mechanisms which influence fatigue life identified. A small reduction in fatigue life was found with decreasing L Irrespective of the position of hold at peak strain in a cycle, the hold time always reduced the fatigue life in comparison with continuously cycled tests at lower strain rates but of equal cycle duration. Tensile holds were found to be most damaging, followed by compression holds. Symmetrical tension-plus-compression holds led to fatigue lives which were very close to those of the continuously cycled tests. In the continuously cycled tests with strain rates down to ~ =6.7x 10 -5 S -1, failure was always transgranular with no indication of creep damage. However, for tests with tensile holds, creep damage was evidenced by grain boundary cavitation and oxidation, which were responsible for a reduction in the fatigue life. The formation of thick oxide scales at the surface observed at longer *Permanent address: Materials Development Laboratory, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, Tamil Nadu, India. 0921-5093/88/$3.50
hold times led to chromium-depleted surface zones in which the carbide precipitates were dissolved. The loss of carbides increased grain boundary sliding which enhanced the formation of grain boundary cracks. These cracks shortened the critical length of surface fatigue cracks which, during fracture of the specimen, linked with the intergranular creep cracks in tensile hold tests. The damaging effect of compression hold was attributed to increased inelastic strain and deformation ratcheting found in the cycle; failure occurred by local accumulation of tensile plastic strain which finally caused tensile necking. 1. Introduction This study was performed as part of the materials programme for the development of the high temperature gas-cooled reactor (HTGR) [1, 2]. Inconel 617 is the principal candidate alloy for heat exchanger components, in which low cycle fatigue (LCF) loadings occur owing to reactor start-up and shut-down procedures and to power transients. Operating temperatures are in the range where creep deformation occurs so that the interaction between L C F and creep is an important aspect of the materials behaviour. At the anticipated operating temperatures of an H T G R , only limited data were available on the L C F behaviour of the candidate material Inconel 6 1 7 t [3-9]. Most of these tests were conducted at tThe German standard designation is Werkstoff Nr. 2.4663 or NiCr22Co12Mo; Inconel 617, the designation given to this material by INCO, the company which introduced the alloy,is used throughout this paper. © Elsevier Sequoia/Printed in The Netherlands
38
strain rates of approximately 10-3 s-l, which is orders of magnitude higher than thermally induced strain rates in the heat-exchanging components. There was therefore a need for fatigue data at lower strain rates especially under creep-fatigue loading and for a full understanding of deformation and fracture mechanisms during high temperature fatigue. Therefore, this investigation was carried out to evaluate the effects of strain rate and hold time on the LCF life of Inconel 617 at 950°C in simulated primary-circuit helium gas environment of an HTGR and to correlate the results with the precipitation behaviour, deformation and damage mechanisms in the material, particularly under various hold time loading conditions.
that the temperature profile remained _+3°C throughout the tests, which were conducted in a pre-evacuated chamber and in flowing impure helium gas of composition specified in Table 2. This composition was agreed on by the German HTGR project partners as the representative environment of the primary circuit. A high resolution gas chromatograph was connected to the input and output of the pressure chamber to monitor the gas composition. The test gas was obtained as a bottled premix. The pressure and flow rate of the gas were 1.8 bar and 3 1 h-1 respectively.
2. Experimentalprocedures i
The Inconel 617 test material was supplied by Vereinigte Deutsche Metallwerke AG., Werdohl, ER.G., under the trade name Nicrofer 5220 Co as hot-rolled and solution-annealed plate 17 mm thick, which had been solution treated for 1 h at 1160°C. The chemical composition is given in Table 1 and the grain size was ASTM 6.5. Specimens were machined to the form shown in Fig. 1 and LCF tests were carried out in the laboratories of Interatom G.m.b.H., D-Bensberg, using MTS servohydraulic test machines of + 1 0 0 kN dynamic load capacity. All experiments were performed under total strain control; the axial strain was measured with an extensometer mounted on the ridges of the test specimen. A nominal strain range of 0.6% and a temperature of 950 °C were employed for all the tests. Uniform heating of the specimen gauge length was accomplished using an induction coil and by carefully monitoring and controlling specimen temperature with Pt-(Pt-Rh) thermocouples fixed in a small hole near the gauge section. Previous calibration with dummy specimens showed
TABLE 1
Impurity p (/~bar)
detail X
_____A Fig. 1. Geometry of the specimen (all dimensions are in millimetres).
Chemical composition of Inconel 617
Element Amount(wt.%)
TABLE 2
I
Ni 54.9
Cr 21.44
Co 11.81
Mo 9.34
AI 1.11
Ti 0.55
C 0.087
Partial pressures of the impurities in the helium gas Hz 500 + 50
CH 4 20 _+5
H20 1.5 -+ 1
CO 15 + 5
CO 2 1.0
N2 5
39
Two types of test were conducted. In the tests for examination of the strain rate effect on fatigue life, a symmetrical continuous triangular waveform was used at strain rates of 2 x 10 -5 s ] to 4 × 10 -3 s-1. In the hold time tests for the study of the creep-fatigue-environment interaction the strain was held constant in a cycle at tensile peak strain, at compressive peak strain, or symmetrically at tensile and compressive peak strains. The peak strain holds were superposed on a triangular wave with g = 4 x 10- 3 s- 1. The waveforms employed are schematically described in Fig. 2. Hold time effects were studied for durations ranging from 5 s to 120 min. In order to facilitate the comparison of fatigue properties, strain rates and hold times have been chosen in such a way that for most of the tests the time for 1 cycle remained equal for different types of loading. For example, the cycle durations using 10 min ten-
CONTINUOUSCYCLING I
|
- -
/x,
|
l
TIME
TENSIONHOLD [
-th
t,)
I
TIME
COMPRESSIONHOLD (to)
i /~
I
sion, using 10 min compression, or using 5 min tension plus 5 min compression hold times and the cycle period for the symmetrical triangular waveform tests with g = 2 x 10-5 s 1 are all equal (Fig. 2). Fracture surfaces were examined by optical and scanning electron microscopy (SEM) to determine the crack initiation and propagation modes. Studies were also conducted on longitudinal sections of failed specimens to characterize the behaviour of surface cracks and the microstructural changes that occurred during fatigue testing. These studies were supplemented by transmission electron microscopy (TEM) investigations for which samples were obtained 1 mm below the fracture surfaces normal to the loading axis. Thin foils were prepared by initially grinding the samples down to 250 /~m, followed by electropolishing in a solution containing 20% H 2 S O 4 and 80% methanol at 25 V and 0°C in a twin-jet apparatus.
/
TIME
ENSION AND COMPRESSIONHOLD ~-- th-~
"Z Z
E
°
/
TIME
] - q c --'~
--cycle period-Fig. 2. Wave shapes employed in LCF testing.
3. Results
3.1. Determination of the number N l of the cycles to macrocrack initiation and number N v of cycles to failure N L and N v were estimated from the plots of absolute peak stress ratio IO'compression/O~ensionl IS. number of cycles. The number of cycles corresponding to a 2% and 50% increase from the plateau value of the stress ratio curve are taken as the life NI to macrocrack initiation and the life Nv to failure respectively. However, the determination of N~ and NF could not be established by this procedure under compression hold conditions because no variation in stress ratio occurred up to the end of specimen life owing to identical stress responses in the tension and compression portions. In this particular case, Nv is taken as the life at rapid onset of tensile stress drop that took place towards the end of the test. 3.2. Influence of test variables on fatigue life A summary of the experimental results showing the influence of strain rate, different types of hold time and different durations of hold time on NI, Nr and time te to failure is presented in Table 3. Figure 3 shows the relationship between the fatigue life reduction factor R and the cycle period for different loading conditions. The fatigue life reduction factor is defined as the ratio of Nv/Nvo, where N v is the life recorded for a
40 TABLE 3
Influence of strain rate, hold condition and hold time on fatigue life at
Specimen number
g in ramp (s- 1)
1694 1437 1724 1727 1729 1454 2131 1699 1450 1712 1449 1698 1445 1713 1707 1708 1716 1715 1717
4.0 × 1 0 - 3 4.0 × 1 0 - 3 4.0 × 1 0 - 4 2.0 × 10 -+ 6.7 x 10 -5 2.0 x 10 -5 4.0 x 10 -3 4.0x 10 3 4.0 x 10 -3 4.0 x 10 -3 4.0 X 1 0 - 3 4.0 x 10 -3 4.0 x 10 -3 4.0 x 10 -3 4.0x 10 3 4.0x 10 3 4.0x 10 3 4.0x 10 3 4 . 0 x l O -3
l
i0 °
'
,
.,f
HoM condition
q
Tension Tension Tension Tension Tension Tension Tension Tension Tension Compression Compression Tension + compression Tension + compression
,
i
, I
w
'
' ' 1
'
,
I,
O tension hold times 13 compression held times
--0
__~__~
..
"n"*,7:22;
V etrl|n
"'"
10 -I
10-1
Fig. 3. The life reduction factor R against the cycle period in minutes for different loading conditions. R is defined as the ratio of N~-/Nvo where N F is the fatigue life recorded for a given strain rate g or wave shape and hold time and Nv0 is the reference life value for continuous cycling with g = 4 × 10-3 s-I (o).
g i v e n s t r a i n r a t e o r h o l d t i m e t e s t a n d Nv0 is t h e r e f e r e n c e life v a l u e f o r c o n t i n u o u s c y c l i n g at g=4×10-3s -1. The results of continuous-cycling tests showed a v e r y g r a d u a l r e d u c t i o n in f a t i g u e life w i t h t h e l o w e r i n g o f t h e i m p o s e d s t r a i n r a t e , i.e. w i t h a n i n c r e a s e in c y c l e d u r a t i o n . A r e d u c t i o n in g b y a f a c t o r o f 2 0 0 c a u s e d a f a t i g u e life r e d u c t i o n o f less t h a n a f a c t o r o f 2.
T =
9 5 0 °C (total strain e t = 0.6%)
HoM time (rain)
Number N I of cycles to crack initiation
Number N~ of cycles to failure
Time to failure (h)
0 0 0 0 0 0 0.083 1.0 1.0 3.0 3.0 10.0 10.0 30.0 120 1.0 10.0 0.5 +0,5 5.0 + 5,0
2530 3000 1900 1610 1775 1250 1080 610 700 335 560 270 300 220 130 -----
3025 3200 2725 2370 2325 1580 1420 830 890 650 740 450 465 380 215 1030 665 2200 1500
2.5 2.7 22.7 39.5 116.3 263.5 3.5 14.5 15.0 32.5 37.0 15.5 77.5 190.0 430.0 17.5 111.5 367.0 251.5
Hold times at the tensile peak strain led to a pronounced reduction in fatigue life. It can be seen that tensile hold periods of only 5 s resulted in shorter fatigue lives. A compressive peak strain hold was also found to cause large reductions in fatigue life but its effect was slightly less than that of tensile hold. Equal hold periods both in tension and in compression yielded reduction factors which were very close to the continuouscycling data with identical cycle duration. It should be noted that fatigue life reduction increased continuously with an increase in hold time irrespective of the type of hold conditions employed in a given cycle; no saturation of the reduction was evident. 3.3. Cyclic stress response as a function of strain rate and type of hold time Figure 4 shows half-fife hysteresis loops for various testing conditions to characterize and compare their effects on material cyclic deformation behaviour. Continuous strain cycling promoted approximately symmetrical stress response in the tension and compression portions of the cycles (Fig. 4(a)). Reducing the strain rate yielded a considerable decrease in the maximum tensile and compressive stress and increased the amount of inelastic deformation in each cycle. Irrespective of the strain rate applied, the stress approached a constant level. This feature has also been found in
41
tensile deformation behaviour of Inconel 617 and other high temperature alloys at 950 °C [10, 11]. At the strain reversals during the continuouscycling experiment with a strain rate of 4 × 10 -3 s -1 (Fig. 4(a) and Table 4) a rapid stress relaxation was observed, indicating an inelastic strain rate faster than the applied total strain rate. At low strain rates, i.e. for g~< 2 × 10 -4 s -~, this rapid relaxation effect was absent because of lower initial stresses. 250 20O
f
f
t 50 I00
~.
50 o
"~ -5o - ~:2xlO-Ss ~
-I00 -150
-200
: 2 x 10-4 s-I
,
-250 (a)
0.1%
250 200 150 I00
o2
5o o
b
-50 -I00 -150 -200 -250 (b)
continuous cycling
IO rnln tenslon hold
5ram tension + 5 rnln compression hole
s "--~
~ _ 0.1%
Fig. 4. Stress-strain hysteresis loops (a) representing the midlife to crack initiation in continuous-cycling tests at different strain rates and (b) taken at mid-life to crack initiation for the continuous-cycling condition ( g = 4 x 10 _4 s - l ) , 10 rain tension hold, 10 rain c o m p r e s s i o n hold and 5 min symmetrical hold.
When hold times were applied, in either tension or compression, a very pronounced stress relaxation occurred (Fig. 4(b)). Even though only short hold times have been applied, the stresses relaxed to nearly zero. In the cycles consisting of tension-only or compression-only hold times, this pronounced stress relaxation yielded asymmetrical hysteresis loops. For tension hold times the inelastic strain in the compressive part of the hysteresis loop was enlarged compared with the inelastic strain in the tensile part and the opposite correlation was found for compression hold times. The cyclic stress response curves for various testing conditions are depicted in Fig. 5. In this figure the tensile stress amplitude is plotted against the fraction N/Nv of fatigue life in order to show more clearly the various deformation stages. The onset of the stress drop towards the end of the test generally indicates the formation of macrocracks and their subsequent growth [12]. Continuous-cycling tests exhibited, after a rapid initial stabilization of peak stresses, a more or less steady state stress response over a large portion of fatigue life (until N / N F <<.0.70). Tests carried out with a 10 min tension hold showed a small rapid initial softening followed by a gradual softening until macroscopic crack growth occurred. It should be pointed out that the relative fraction of fatigue life spent in lhe crack propagation stage was considerably longer in the 10 min tensile hold condition than in the continuous-cycling test with g = 4 x 10 3 s- l Compression holds alone and tension-pluscompression holds initially caused slight cyclic strain hardening followed by softening. Symmetrical holds developed slightly higher tensile peak stress than the holds introduced in compression or tension alone.
TABLE 4 Computed values of elastic strain Aeel , plastic strain Aep, rapid relaxation strain epr, and effective plastic strain Aep' as a function of strain rate
Specimen number 1694 1724 1727 1729 1454
g (s 1)
4x10 3 4 x 10 -~ 2 x 10 4 6.7 x 10 5 2.0 x l O- 5
Aet (%)
0.6 0.6 0.6 0.6 0.6
Aeel (%)
0.21 0.17 0.16 0.13 0.09
Ae (%3
0.33 0.39 0.43 0.47 (I.49
Rapid relaxation strain 6pr (%)
Effective plastic strain
In tension
In compression
Aep' = Aep + 2e F
0.03 0.015 ----
0.03 0.015 ----
0.39 0.42 0.43 0.47 0.49
Ae~ is the total strain range, A e p = A e , - 2 d I E and %r = -IAddl/E w h e r e E is Young's m o d u l u s , d the stress amplitude in the cycle and A d d the stress drop at strain reversal.
42
• 4. I0 3 s 4
1
~lOO 0 tOmin tension hold
200
leo
O 200
I0 rain compressionhold.
;,--____._.._
100 0 200
5min tension +
leo
0 200
Fig. 6. Photograph of broken specimens tested at g= 4 × 10-3 s t under the followingconditions: (a) continuous cycling;(b) 10 min tension hold; (c) 10 min compression hold; (d) 10 min tension-plus-compressionhold. Geometrical instability in the gauge portion of the samples subjected to tension-onlyand compression-onlyholds can be clearlyseen.
k
~= Z' 10.4 S"1 10l O 20| E=2.1OSs "1
Alloy 800H in hold time tests conducted at
I0|
850 °c [8]. 0
0,1
i
i
i
I
,
i
i
,
0,2
0,3
0.4
0.5
0.6
0.7
0.8
0.9
- -
I,O
N/NF
Fig. 5. Cyclic tensile stress response curves for various loading conditions studied. The tensile stress amplitude is plotted against the fraction N/Nv of fatigue life. The total strain range Ae is 0.6% throughout. By decreasing the strain rate from 4 x 10-3 to 2 x 10 -5 s -], the value of the saturation stress was reduced from 228 to 64 MPa. Such a strong g dependence of saturation stress was also noted in the LCF tests conducted in air at 850°C [13]. In hot tensile tests with Inconel 617 at 950 °C the ultimate tensile strength shows the same dependence on strain rate as the saturation stress during the L C F tests [11, 14]. 3.4. Fractographic observations The effect of loading conditions on the physical appearance of the failed specimens is shown in Fig. 6. There was no significant departure from the original geometry after L C F testing with continuous-cycling waveform (Fig. 6(a)) or symmetrical hold (Fig. 6(d)). It can be seen that tensile hold (Fig. 6(b)) induced geometrical instability into the gauge length in the form of bulging whereas in compression hold (Fig. 6(c)) the instability was manifested in the form of necking. Similar observations have been made on Inconel
In the continuous-cycling tests a clear transition in fatigue crack initiation behaviour was found from transgranular stage I shear cracking at g > 2 x 10- 4 s- t to intergranular initiation at g < 6 . 7 x 10 -5 s -1. Crack propagation was found to be transgranular stage II at all the strain rates in continuous cycling. In the 1 min tension hold test, cracks initiated transgranularly by stage I shear cracking on the specimen surface (Fig. 7(a)) whereas crack propagation occurred by mixed mode (Fig. 7(b)). In the test with tensile holds of greater than 10 min, crack initiation occurred intergranularly (Fig. 7(c)). The fracture surface of the samples with tensile hold periods of longer than 1 rain contained a fairly large number of grain boundary cracks in the interior of the specimens. The compression hold test specimens displayed a dimple fracture (Fig. 7(d)) similar to a tensile fracture accompanied by necking. From the small remaining fracture surface in compression hold tests the crack initiation mechanism could not be deduced. The fracture surface of symmetrical hold tests showed transgranular stage I crack initiation and mixed-mode propagation. 3.5. Precipitation behaviour Microstructural features such as carbide precipitation can play an important role in cyclic
43 stress-strain behaviour, the modes of crack initiation and propagation, and fatigue ductility. In the present study, this aspect was investigated using electron diffraction, energy-dispersive X-ray analysis (EDXA) and SEM backscattered electron image techniques.
3.5.1. Primary carbides Two types of carbide not dissolved by the solution treatment were found in all the samples studied. Relatively coarse cube-shaped particles were identified as primary Ti(C, N) particles by EDXA. The second type of particle appeared as stringers in both intragranular and intergranular regions (Fig. 8). An E D X A of these particles indicated that they contained both chromium-rich and molybdenum-rich regions corresponding to M 2 3 C 6 and M 6 C types of carbide, which have been identified as the grey and white areas respectively. The proximity of these phases confirms that M6C is a result of M23C~ transformation [15, 16].
Fig. 7. (a) Fracture surface of a 1 min tensile hold sample, transgranular stage I crack initiation and stage II propagation; (b) mixed-mode propagation in the interior of a 1 min tensile hold specimen; (c) intergranular crack initiation in a 120 min tensile hold sample; (d) fracture surface showing dimple fracture in a 10 min compression hold test.
3.5.2. Development of secondary carbides as a function of loading condition The dark field TEM micrograph shown in Fig. 9(a) illustrates the morphology and distribution of M23C6 precipitated along the grain boundary in the samples tested at a strain rate g of 4 x 10 -3 s-1 with zero hold. These carbides had an f.c.c, ordered crystal structure coherent with one of the neighbouring grains, as shown by electron diffraction. One of the very fine intergranular M23C6 particles is shown in Fig. 9(b). No secondary M~C was found in samples continuously cycled at g= 4 x 10 -3 s-l; this is consistent with the prediction of an incubation time for its appearance at 950°C [16]. The specimen life t~of 2.7 h for g= 4 x 10 -3 s- ~is far less than the incubation time reported [16] for onset of M~C precipitation. At g < 4.0 x 10 - 4 s- ~, copigus amounts of M~C precipitated together with M23C6 were found in both intragranular and intergranular regions (Fig. 9(c)). The introduction of a 1 min tensile hold at the peak strain led to an increase in the volume fraction and size of intragranular M23C6 (Fig. 10(a)). An increase in both the size and the density of grain boundary carbides was also noted. The carbides formed along the grain boundary were predominantly M23C6, a small amount of M 6 C w a s also detected. In addition to the grain boundaries,
44
~
! 20 Dim
II
Fig. 8. SEM image obtained in backscattered electron mode showing undissolved massive precipitates in intragranular regions. White precipitates correspond to MaC and grey precipitates to M23C 6 (T=950°C; ramp rate ~=4x 10 3 s-~; tensile hold for 10 min). deformation bands and twin boundaries were also preferential sites for carbide precipitation (Fig. 10(b)). Secondary precipitation of finely dispersed Ti(C, N) was detected in tests with tensile holds of 1 rain or longer in the intragranular regions. A further increase in tensile hold to 10 min resulted in the rapid coarsening of grain broundary precipitates and increased interparticle spacings (Fig. 10(c)). Raising the duration of hold to 120 min did not produce any further significant changes in the grain boundary precipitate size, density, morphology and interparticle spacings. The SEM backscattered electron image in Fig. 1 l(a) illustrates the typical microstructure obtained in samples tested with a 5 min tension plus 5 min compression hold. It can be seen that discrete and large precipitates of M6C (white phase) and M23C 6 (grey phase) coexisted on large-angle grain boundaries, but few were found in the matrix. The proportion of M6C on grain boundaries in the 5 min symmetrical hold test was very much higher than that observed in either the longest-duration tensile hold or the 10 min compression hold test.
Fig. 9. (a) Dark field TEM m]crograph illustrating grain
boundary M23C6 (T=950°C; g=4xl0 -3 s-I); (b) intra-
granular M23C6 carbides with interracial dislocations (T= 950°C; g= 4 × 10 -3 S-I); (C) backscattered SEM image showing Ti(C, N) (dark phase), M23C6 (greyphase) and M6C (white phase) particles (T= 950°C; g= 2 x 10 -5 s- i). In the 10 min compression hold test a nearly continuous network of grain boundary carbides was observed, as shown in Fig. 1 l(b). In the tests with holds of longer duration, intragranular M23C6 and M6C acquired a spherical shape (Fig. 1 l(c)) compared with the cuboidal shape which was commonly observed in continuous-cycling tests or tests with a short hold time. The formation of )/precipitates (Ni3(A1, Ti)) in Inconel 617 has been described in the literature [16, 17]. In the present study, the introduction of
45
~f
E ¸¸¸
Fig. 10. (a) TEM micrograph showing intragranular M23C6 particles in the samples tested with a 1 min tension hold; (b) secondary-electron SEM image showing precipitation on deformation bands in the samples tested with a 1 min tension hold; (c) intragranular and intergranular precipitate distribution in samples tested with a 10 rain tension hold.
a compression hold alone or in combination with a tensile hold induced serpentine-shaped ~' phase identified by diffraction, but the amount detected was so small that no effect on ductility and fatigue life is expected.
3.6. Deformation behaviour Regardless of the loading condition, the alloy exhibited generally subgrain formation in the
Fig. 11. (a) SEM backscattered electron image illustrating M23C6 (grey phase) and M6C (white phase) precipitates in the samples tested with a 5 min symmetrical hold; (b) secondaryelectron SEM image showing continuous network of grain boundary carbides in the samples tested with a 10 min compression hold; (c) bright field TEM micrograph of spherical M23C6 precipitates in the samples subjected to a l0 min compression on hold (the precipitates are found to be spherical at longer hold times in contrast with cube-shaped particles seen at shorter holds).
majority of the grains (Fig. 12(a)). Multiple slip in some of the grains was also observed, suggesting that the mode of deformation was predominantly homogeneous at 950 °C. In the tests with a tensile hold, the alloy also showed some tendency to form dislocation tangles in the matrix and dislocation pile-ups at the grain boundaries (Fig. 12(b)). However, in spite of dislocation pile-ups, no intergranular cracking could be observed. Under a 5 min symmetrical hold, the substructure
46
Fig. 13. Transgranular secondary crack in the specimen tested with g= 4 x 1 0 - 3 s - 1 . The arrow in the top right-hand corner indicates the loading direction. 4. Discussion
4.1. Influence o f the strain rate on fatigue life reduction
Fig. 12. Bright field TEM micrographs showing(a) substructure in samples tested with g= 4 x 10 3 s- ', (b) the pile-up of dislocations on grain boundaries and dislocation tangles in the intragranular regions and (c) interaction between coherent M23C 6 and dislocations (carbide precipitates have been identifiedthrough moire fringe contrast). indicated extensive interactions between coherent precipitates and dislocations (Fig. 12(c)). In this figure the coherent M23C6 precipitates were identified by moire fringe contrast. It is reasonable to correlate such interactions with the slight cyclic hardening observed for this loading condition.
It is evident from Fig. 3 that decreasing the strain rate results in a small reduction in fatigue life. Associated with this reduction are a decrease in stress amplitude, an increase in inelastic strain range (i.e. an increase in the width of the hysteresis loop at zero stress; see Fig. 4 and Table 4) and a change in crack initiation mode from transgranular at g >t 2 x 10-4 s- 1 to intergranular at g ~ < 6 . 7 x l 0 - S s -1. At the lowest strain rate, cracks were initiated by the oxidation of surfaceconnected grain boundaries. Intergranular cracks propagated to only one grain diameter before a transition to transgranular propagation took place. From the studies of secondary cracks in longitudinal sections, it was found that the transgranular crack propagation path remained sharp and narrow (Fig. 13). This observation suggests very little interference of the oxidation process in the crack propagation mechanism. Furthermore, microstructural investigations revealed no evidence of any creep damage in the bulk even at the lowest strain rate investigated in continuouscycling tests. On the basis of detailed investigations of the influence of strain rate at 850 °C in air, it has been argued that the reduction in fatigue life with decreasing strain rate results from the combined effects of increased inelastic strain and from oxidation of surface connected grain boundaries and environmentally accelerated mixed-mode propagation [13]. In the range of strain rates examined in this investigation, the contribution from environmental attack to reduction in life appeared only at g < 6 . 7 x 10 -5 s -1 in the form of oxida-
47
tion-assisted intergranular crack initiation. Moreover, the reduction in fatigue life occurred with decreasing g, even in the range of strain rates (~= 2 x 10-4-4 x 10 -3 s -1) where entirely transgranular crack initiation and growth prevailed. These observations demonstrate that fatigue life is largely governed by the increase in effective plastic strain component of the cycle as g is lowered. 4.2. Hold time effects 4.2.1. Mean stress under hold time conditions From Fig. 4(b), it can be observed that the addition of either tension-only or compressiononly holds causes no significant mean stress in the hysteresis loops. The capacity of the material to withstand the mean stress depends on the ratio of inelastic to elastic strains in a cycle [18]. If the inelastic strain is large relative to the elastic strain, as in the case of the present experiments, the mean stress which can develop is very low and consequently mean stresses do not influence cyclic life significantly. 4.2.2. Relaxation The damage behaviour during the hold period could be best described by characterizing the stress relaxation behaviour as a function of time. A typical mid-life stress relaxation response curve for the 120 min tensile hold test is shown in Fig. 14. It can be seen that rapid relaxation down to half of the maximum stress occurred in less than 1 s. Further, the inelastic strain rate gr associated with relaxation strain continuously decreased with increasing hold time from approximately 6.5 x 10 4 t o 4 × 10 9 s - l .
In Fig. 15 the inelastic strain rates associated with stress relaxation for different hold conditions were compared with minimum-creep-rate data available from creep tests for various stress levels on the same heat. The strain rates during stress relaxation are generally higher than the corresponding minimum creep rates at all the stress levels. The shape of relaxation rate curves resemble the shape of the minimum creep rate vs. stress curves at a lower strain rate but distinctly deviate at higher rates. Inelastic strain rates found in the rapid relaxation period correspond to those which are expected for precipitation-free matrix deformation, while those observed in slow relaxation period are typical for creep deformation. 4.2.3. Cavitation damage in tensile hold tests The build-up of tensile inelastic relaxation strain caused severe cavitation damage within the bulk of the material. Figure 16(a) clearly illustrates the development of intergranular cavities (R type) associated with grain boundary precipitates adjacent to the fracture surface in a 1 min tension hold test. Figures 16(b) and 16(c) describe the cavity formation and cavity linkage on grain boundaries in the samples cycled with 10 and 1 20 rain tension holds in the regions remote from the fracture surface. R-type cavities form at grain boundaries and second-phase particles.
n0, s0 100 lu e00 ~s0 ~ ~ ,~ ' , ' ~ , ' i " ~ ",0~,~',',?' /~'//~-'~--"~i ~ v~ llimin l~.~n.~u |~ tension
!
////°
,.o
........////j
. .........
............. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Ill
....
;,,,j,
....
no
0
o lo'
ao'
m'
lo't/sec~
~o'
Fig. 14. Mid-life stress relaxation curve in a 120 rain tensile hold test (it, inelastic relaxation strain rate),
~ .... 50
i .... IH
I .... 150
i
. . . .
ZooG./Mra25o
Fig. 15. Relaxation strain rate data for different hold conditions. M i n i m u m - c r e e p - r a t e data from stress rupture tests are also included in the figure for c o m p a r i s o n wilh relaxation strain rate data.
48
Fig. 17. (a) Grain boundary cracks originating from wedge cracks in the near-surface regions free from carbide precipitation (10 min tension hold); (b) thick oxide layer formed on the surface (the crack faces of secondary cracks are also oxidized, and internal oxidation ahead of the crack tip can also be seen) ( 120 min tensile hold).
Fig. 16. (a) Intergranular cavitation damage adjacent to the fracture surface in the sample tested with a 1 min tension hold (cavities are associated with grain boundary precipitates); (b) cavity coalescence on grain boundaries oriented at 90 ° to the loading axis (10 min tension hold); (c) intergranular cavity damage in the samples subjected to a 120 min tension hold.
In addition to the bulk cavity damage, a large number of grain boundary cracks were observed in the surface regions particularly when the tension hold times were more than 10 min (Fig.
17(a)). A closer examination showed that these cracks formed in the near-surface regions which were free from M23C6 precipitation nucleate as wedge-type cracks. This observation supports the conclusion that sliding is higher at grain boundaries free from carbides [19] which facilitates linking of cavities to form grain boundary cracks. Furthermore, the damage due to grain boundary oxidation was found to be considerable at longer tensile holds. Figure 17(b) shows the surface oxide layer and oxide-filled subcracks in the failed sample with an imposed tensile strain hold of 120 min. A n E D X A of the oxide layers demonstrated that these were basically of chromia. From Fig. 17(b), it is also apparent that the oxygen diffusion ahead of the crack caused internal oxidation of aluminium. A zone of 10/~m from the surface was also found to contain discrete particles of aluminium oxide. Since most of the oxides are intrinsically brittle, the fatigue
49
cracks in tests with longer hold periods initiated in oxides which are formed preferentially at surface-connected grain boundaries. The oxidation-induced surface intergranular cracks penetrated deeply into the interior and merged with independently formed intergranular wedge cracks in the near-surface regions and R-type cavities in the bulk. 4.2.4. Compression hold tests Compression hold tests did not exhibit creep cavities in the specimen interior. In addition, none of these features such as a thick surface oxide layer, carbide-free zones and wedge cracking in the near-surface zones which have been prominently displayed in long period tensile hold tests was found in compression hold tests (Fig. 18). Bulk creep damage is unlikely to develop during the periods of compressive stress relaxation. Furthermore, compressive holds have been generally acknowledged as capable of sintering creep cavitation damage. Therefore the damaging effect of compression hold on life reduction could only be ascribed to the increased amount of inelastic strain in a cycle resulting mainly from stress relaxation. It must be remembered that, during compression holds, test specimens exhibited geometrical instability (Fig. 6) and produced tensile necking failure. This has occurred in the absence of significant tensile mean stress. The reason why compression hold strain cycling leads to tensile fracture has been interpreted by Manson [20]. The progressive accumulation of tensile inelastic strain with continued cycling in the tensile half and its detrimental effect compared with the accumulation of relaxation creep strain during compression hold leads to tensile fracture.
4.2.5. Symmetrical hold tests The optical micrographs presented in Fig. 19 illustrate the results for 5 min symmetrical hold tests. In contrast with the uniformly distributed intergranular cavitation damage in the specimen interiors of tensile hold tests, the cavity damage has been confined close to the fracture surface (Fig. 19(a)). No wedge cracks were seen in the carbide-depleted zone in the near-surface regions (Fig. 19(b)) as opposed to the observations in tensile hold tests. The few intergranular cracks which have merged with surface cracks appeared to be a consequence of internal oxidation of the grain boundaries. The non-occurrence of wedge cracking in the carbide-depleted region would increase the critical length of the fatigue crack before the onset of interaction with the cavity damage in the central region of the specimen is established. This interaction and the increased amount of inelastic strain in the cycle are responsible for the slightly inferior fatigue resistance of the symmetrical hold tests compared with continuous-cycling no-hold tests of equal cycle duration. In spite of the higher amount of inelastic
'o.T: /~
....
•
~. . . . . . . . ,
-
-".~,
~
.'" : "~'" , , ' ~ . 1 / ~ ' ; )
.~-
" ,:
:
" "
~ 1 5 0 p m
Fig. 18. View of the surface and subsurface regions in the longitudinal section of the samples tested with a 10 rain compression hold.
. . . . . . . . . . .
'
'
.................
:
150
Fig. 19. (a) Ca~tation damage near the fracture surface in the 5 min symmetrical hold test; (b) carbide-depleted zone in the surface regions and secondary cracks in this zone.
50 strain in a cycle and fatigue-cavity damage interaction, symmetrical holds exhibited better fatigue resistance than compression-only holds probably because of the non-existence of deformation ratcheting in the cycle. It has been suggested that, if compressive hold periods follow the tensile hold periods, sintering of cavities formed during the tensile hold period takes place during the compressive hold [21[. In fact, the experimental findings of Majumdar and Maiya [22] on AISI 304 stainless steel showed that symmetrical holds produce only transgranular failure contrary to our observations of an interaction between the fatigue and cavity damage. It should be noted from Fig. 19(b) that the cavities became rather large, corresponding to the size of the grains, and were definitely not associated with grain boundary carbides as in the case of tensile-only holds. Furthermore they were oriented at about 45 ° to the stress axis. These observations indicate that these cavities nucleate and grow by irreversible shear deformation.
4.3. Corrosion effects In this investigation the main corrosion effects observed in simulated primary-circuit helium of the HTGR (Table 2) were the formation of predominantly chromium-based oxide scales at the surface, internal oxidation of aluminium beneath the surface oxide layer, in the adjoining areas of crack faces and at the crack tip, and the carbidefree zone in the near-surface region. These effects are consistent with the data reported on isothermal corrosion behaviour of this alloy at 950 °C in an HTGR helium environment [23, 24]. In the simulated HTGR helium environment (Pra~o above 0.5 pbar) the aluminium content of Inconel 617 is too low for a surface scale based on AIzO 3 to form; A1203 exists only as internal oxide beneath the stable chromium oxide layer [24]. The carbide-free zone in the near-surface region results partially from the depletion of chromium which is consumed in oxide formation. An E D X A of the carbide-free area has shown a chromium-depleted zone of only 30 pm, which is in agreement with the data reported on isothermal exposures in HTGR helium at 950 °C [24]. Because the carbide-free area is much larger than the chromium-depleted zone, it is not reasonable to assume that the depletion of chromium is entirely responsible for the development of carbide-free zones. This indicates that decarburization was occurring in the near-surface regions. In
fact, the analysis of the test atmosphere of the outlet of the LCF test chamber by gas chromatography indicated a considerable increase in CO content in the longer-hold-time tests. Decarburization has also been noted for this alloy after stress-free exposure in HTGR helium atmospheres at 950 °C [23]. 5. Conclusions
(1) Decreasing strain rate caused a reduction in the number of cycles to fatigue crack initiation and to failure. In the range of strain rates examined ( g = 4 x 10-3-2 x 10 -5 s-l), the major factor responsible for fatigue life reduction was the continuous increase in inelastic strain component in a cycle as the strain rate was lowered. The contribution from environmental attack was found only at g<6.7 x 10 -5 s -1 in the form of oxidation-assisted intergrannlar crack initiation. (2) Irrespective of the position of hold at peak strain in a cycle, hold time always reduced fatigue life compared with continuous-cycling tests. (3) Hold periods in tension led to significant reductions in fatigue life; these reductions were slightly larger than those observed for compression-only hold tests and considerably larger than those involving symmetrical hold periods in both tension and compression. (4) Hold periods in compression produced failure by necking. Owing to the large amount of relaxation, an unbalanced tensile inelastic strain occurred and was accumulated during cyclic loading. No cavity damage was observed. (5) The reduction in fatigue life which occurred when tensile hold periods were applied was principally due to the interaction between the surface-initiated fatigue crack and interior grain boundary creep cavitation damage formed during periods of tensile stress relaxation. (6) At longer tensile holds, a large number of grain boundary cracks were found in the nearsurface regions depleted of grain boundary carbides. Cracking in these regions is believed to shorten the critical length of surface fatigue cracks which establishes the interaction with bulk cavity damage associated with the carbides. (7) Oxidation played a dominant role in fatigue initiation and early propagation at longer tensile holds. There was a progressive change in crack initiation mode from transgranular stage I at short tensile hold periods to intergranular
51
initiation at tensile hold periods longer than 10 min. (8) The corrosion effects observed in the HTGR helium atmosphere consisted of the formation of predominantly chromium-based oxide scales at the surface, internal oxidation of the aluminium beneath the surface oxide layer, in the adjoining areas of crack faces and ahead of the crack tip, and formation of carbide-free zones in the near-surface regions. The carbide-free zones formed as a result of consumption of chromium by oxide layer formation and because of decarburization occurring in the HTGR helium atmosphere. Acknowledgments The authors are grateful for the experimental work done by the staff of the Fatigue Laboratory of Interatom G.m.b.H. and the support by the Metallography and TEM Laboratories at the Department for Reactor Materials, Nuclear Research Centre, Jiilich. The careful reading of the manuscript by Mr. E J. Ennis is gratefully acknowledged. K. B. S. Rao is grateful to Shri C. V. Sundaram, Dr. Placid Rodriguez and Dr. S. L. Mannan, all of Indira Gandhi Centre for Atomic Research, Kalpakkam, India, for support during the period of his stay at Jfilich under the Indo-German bilateral agreement. The work has been performed within the framework of the German HTGR project Prototype Plant for Nuclear Process Heat and has been sponsored by the State of Northrhine Westfalia and the Minister for Research and Technology of the F.R.G. References 1 H. Nickel, F. Schubert and H. Schuster, Proc. Conf. on Gas-cooled Reactors Today, Bristol, September 10-24, 1982, BNES, London, pp. 173-178. 2 Status of metallic materials development for application in advanced high-temperature gas-cooled reactors, Nucl. Technol., 66(1984) 11-721. 3 H. E Meurer, H. Breitling and E. D. Grosset, Proc. Conf. on the Behaviour of High Temperature Alloys in Aggressive Environments, Petten, October 15-18, 1979, Metals Society, London, i980, pp. 1005-1016. 4 H. E Meurer, H. Breitling and E. D. Grosser, Proc.
Specialists' Meet. on High Temperature Metallic Materials for Application in Gas Cooled Reactors, in IAEA Rep. 158, May 1981, pp. U 1 - U l l (International Atomic Energy Agency). 5 J.P. Strizak, C. R. Brinkman and E L. Rittenhouse, Proc. Specialists" Meet. on High Temperature Metallic Materials for Application in Gas Cooled Reactors, in IAEA Rep. 158, May 1981, pp. T1-T15 (International Atomic Energy Agency). 6 H. Hattori, M. Kitagawa and A. Ohtomo, J. Soc. Mater. Sci., Tokyo, 32(1983)667-671. 7 M. A. Burke and C. G. Beck, Metall. Trans. A, 15 (1984) 661-670. 8 H. E Meurer, G. K. H. Gnirss, W. Mergler, G. Raule, H. Schuster and G. Ullrich, Nucl. Technol., 66 (1984) 315-323. 9 K. Bhanu Sankara Rao, H. Schiffers and H. Schuster, P~oc. Int. Conf. on High Temperature Alloys -- Their Exploitable Potential, Petten, October 15-17, 198.5, Elsevier, Amsterdam, 1987, pp. 411-422. 10 U. Bruch, E J. Ennis, E. teHeesen, Nucl. Technol., 66 (1984) 357-362. 11 A. Abdel Azim, Jiil-Rep., to be published (Nuclear Research Centre JiJlich). 12 K. Bhanu Sankara Rao, M. Valsan, R. Sandhya, S. K. Ray, S. L. Mannan and E Rodriguez, Int. J. Fatigue, 7(1985) 141-147. 13 K. Bhanu Sankara Rao, H. Schiffers, H. Schuster and H. Nickel, Metall. Trans. A, 19 (1988) 359-371. 14 H. P. Meurer, Interatom G.m.b.H., unpublished research work, 1984. 15 R. F. Decker and C. T. Sims, in C. T. Sims and W. C. Hagel (eds.), The Superalloys, Wiley, New York, 1972, pp. 33-77. 16 H. Kirchh6fer, J. Rottmann, E Schubert and H. Nickel, Jiil-Rep. 1903, March 1984 (Nuclear Research Centre Jiilich). 17 P. P. Schepp, Doctoral Thesis, Universit~t D-ErlangenNiirnberg, 1983. 18 D. C. Lord and U F. Coffin, Jr., Metall. Trans., 4 (1973) 1647-1653. 19 R. yon der Gracht, P. J. Ennis, A. Czyrska-Filemonowicz and H. Schuster, Proc. Int. Conf. on Creep, Tokyo, April 1986, Japan Society of Mechanical Engineers, Tokyo, 1986, pp. 123-128. 20 S. Manson, in A. E. Carder, A. J. McEvily and C. H. Wells (eds.), Fatigue of Elevated Temperatures, ASTM Spec. Tech. Publ. 520, 1973, pp. 744-782. 21 J. B. Conway, R. H. Stenz and J. T. Berling, Fatigue, tensile and relaxation behaviour of strainless steels, Publ. TID 26135, 1975 (National Technical Information Centre, U.S. Department of Commerce). 22 S. Majumdar and P. S. Maiya, Proc. 2nd Int. Conf. on the Mechanical Behaviour of Materials, Boston, MA, August 16-22, 1976, American Society for Metals, Metals Park, OH, 1976, pp. 921-928. 23 W.J. Quadakkers, H. Schuster, Werkst. Korros., 36 ( 1985) 141-150. 24 W.J. Quadakkers, Werkst. Korros., 36(1985) 335-347.